thesis corrected approved.pdf

Production of inorganic nanohybrids by the

templating of carbon and peptide nanostructures

A thesis submitted to The University of Manchester for the degree of

Doctor of Philosophy

In the Faculty of Engineering and Physical Sciences

2013

Yanning Li

School of Materials

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Table of Contents

List of tables…………………………………………………………………..…..8

List of figures………………………………………………………………..……9

List of abbreviations…………………………………………………………..…26

List of symbols………………………………………………………………..…29

Abstract…………………………………………………………………….……30

Declaration………………………………………………………………………31

Copyright………………………………………………………………………..32

Acknowledgements…………………………………………………………..…33

Chapter 1 Introduction…………………………………………………………..34

1.1 Overview ……………………………………………………………………34

1.2 Aims…………………………………………………………………………35

1.3 References……………………………………………………………...……37

Chapter 2 Literature Review……………………………………………….……40

2.1 Sol-gel chemistry……………………………………………………………40

2.2 CNT-Inorganic nanohybrids…………………………………………………43

2.2.1 Introduction to carbon nanotubes………………………………...……43

2.2.1.1 Structures………………………………………………….………43

2.2.1.2 Properties……………………………………………….…………44

2.2.1.3 Synthesis………………………………………………….………46

2.2.1.4 Applications………………………………………………………46

2.2.2 Functionalization of CNTs……………………………………….……47

2.2.2.1 Covalent functionalization………………………………..………48

2.2.2.2 Non-covalent functionalization…………………………...………51

2.2.3 CNT-inorganic nanohybrids ……………………………………….…57

2.2.3.1 Synthesis……………………………………………………….…58

2.2.3.2 CNT-SiO2 hybrids…………………………………………..……59

2.2.3.3 CNT-TiO2 hybrids………………………………………..………63

2.2.3.4 Inorganic nanotubes………………………………………………69

2.3 Peptide self-assembly and mineralization…………………………...………72

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2.3.1 Introduction……………………………………………………………72

2.3.2 Strategy for peptide self-assembly……………………………….……72

2.3.2.1 β-sheets and α helices………………………………………….…72

2.3.2.2 Peptide amphiphiles………………………………………………75

2.3.2.2.1 All-amino acid peptide amphiphiles……………………...……75

2.3.2.2.2 Lipidated peptides……………………………………………78

2.3.2.3 Aromatic short peptide derivatives………………………..………79

2.3.3 Controlled self-assembly of peptides…………………………….……81

2.3.3.1 PH/ionic strength triggered ………………………………………81

2.3.3.2 Enzyme triggered …………………………………………...……82

2.3.4 Mineralisation…………………………………………………….……84

2.3.4.1 Biomineralization…………………………………………………84

2.3.4.2 Biomimetic mineralization………………………………..………88

2.4 Graphene and graphene based nanocomposites……………………..………91

2.4.1 Introduction to graphene………………………………………………91

2.4.1.1 Structure and properties of graphene……………………..………91

2.4.1.2 Production of graphene…………………………………...………94

2.4.1.2.1 Micromechanical cleavage…………………………...………95

2.4.1.2.2 Liquid phase exfoliation………………………………..……95

2.4.2 Graphene based nanocomposites and nanohybrids ………………….101

2.5 References…………………………………………………………….……105

Chapter 3 Experimental Methods……………………………………….…..…127

3.1 Materials……………………………………………………………………127

3.2 Experimental procedure……………………………………………………127

3.2.1 Synthesis of alignedCNT arrays by injection CVD method…………127

3.2.2 Adsorption Study of the surfactants on CNTs………………….……128

3.2.2.1 Adsorption of the surfactants on aligned CNT arrays………..…128

3.2.2.2 Adsorption of the surfactants on randomly aligned CNT

networks…………………………………………………………………134

3.2.2.3 Desorption of the surfactants from CNT arrays in H2O…………135

3.2.2.4 Freundlich adsorption isotherm …………………………....……135

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3.2.2.5 Competitive binding from the Fmoc-AAs library on graphite…..136

3.2.2.6 Switchable surface chemistry……………………………..……..137

3.2.3 Synthesis of CNT-inorganic nanohybrids……………………………138

3.2.3.1 Synthesis of silica coated Fmoc-AA functionalized CNTs…..…138

3.2.3.2 Synthesis of TiO2 coated Fmoc-AA functionalized CNTs……...140

3.2.3.3 Combined sites…………………………………………….……142

3.2.4 Graphene and graphene based nanocomposites and nanohybrids..…143

3.2.4.1 GO-Inorganic nanohybrids ………………………………….…143

3.2.4.1.1 Preparation of aqueous dispersion of GO……………….…143

3.2.4.1.2 Preparation of GO-TiO2 nanohybrids………………………144

3.2.4.1.3 Preparation of GO-SiO2 nanohybrids………………………144

3.2.4.2 bwGO-Inorganic nanohybrids …………………………..………145

3.2.4.2.1 Preparation of bwGO dispersion……………………………145

3.2.4.2.2 Synthesis of bwGO-TiO2 nanohybrids…………………...…145

3.2.4.3 Exfoliated graphene (EG)-Inorganic nanohybrids………………146

3.2.4.3.1 Preparation of graphene dispersion ……………………...…146

3.2.4.3.2 Preparation of EG-TiO2 nanocomposites and nanohybrids…147

3.2.5 Mineralization of peptide self-assembled hydrogels…………………148

3.2.5.1 Fmoc-Y hydrogel preparation………………………………...…148

3.2.5.2 Fmoc-FY hydrogel preparation …………………………………148

3.2.5.3 Characterization …………………………………………………148

3.2.5.4 Silicification of Fmoc-Y gel ………………………………….…149

3.3 Analytical techniques………………………………………………………150

3.3.1 Scanning Electron Microscopy (SEM)………………………………150

3.3.2 Transmission Electron Microscopy (TEM)………….………………150

3.3.3 Energy Dispersive X-ray Spectroscopy (EDX)………………………153

3.3.4 Reversed-phase high-performance liquid chromatography (RP-HPLC)...153

3.3.5 Contact angle measurement………………………………………..…154

3.3.6 Raman spectroscopy …………………………………………………156

3.3.6.1 Background…………………………………………………...…156

3.3.6.2 Raman characterization of the exfoliated samples………………158

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3.3.7 Atomic Force Microscopy (AFM)……………………………………159

3.4 References …………………………………………………………………161

Chapter 4 Dynamic Interaction of Fmoc-AAs with CNTs…………………..…163

4.1 Introduction ……………………………………………………………..…163

4.2 Synthesis of aligned MWNT arrays by injection CVD method ………...…163

4.3. Interaction of surface modifiers with CNTs……………………………….166

4.3.1 Adsorption behavior of modifiers on CNT aligned CNT arrays….…166

4.3.2 Adsorption behavior of modifiers on randomly oriented CNT

Networks…………………………………………………………………170

4.3.3 Desorption behavior of the modifiers in excess of water……………171

4.3.4 Freundlich isotherm model …………………………………………172

4.3.5 Competitive binding from the Fmoc-AAs library on graphite………173

4.3.6 Switchable surface chemistry…………………………………..……176

4.4 Conclusion…………………………………………………………………178

4.5 References…………………………………………………………………180

Chapter 5 Synthesis of CNT-inorganic nanohybrids and the corresponding

inorganic NTs using Fmoc-AAs as surface modifier………………………….181

5.1 Introduction…………………………………………………………..……181

5.2. Synthesis of CNT-silica nanohybrids using Fmoc-AAs as surface

modifier………………………………………………………………….182

5.2.1 Synthesis and morphology characterization……………………….…182

5.2.2 Discussion on the role of Fmoc-AA functionalization in controlling the

morphology of the hybrids……………………………………..……189

5.2.3 Growing mechanism of silica coating on Fmoc-AA functionalized

CNTs……………………………………………………………………..…189

5.2.4 Kinetics for silica growth………………………………………….....191

5.2.5 Annealing………………………………………………………….…193

5.3 Synthesis of CNT-TiO2 nanohybrids using Fmoc-AAs as surface

modifier………………………………………………………………….…194

5.3.1 Synthesis and morphology characterization ……………………...…194

5.3.2 Mechanism for the formation of TiO2 coating on the functionalized

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CNTs………………………………………………………………………...…196

5.3.3Effect of CNT to TBOT ratio on the hybrid morphology……………198

5.3.4 Effect of modifier to CNT ratio on the hybrid morphology…………201

5.3.5Kinetics for TiO2 growth …………………………………………..…202

5.3.6 Synthesis of TiO2 NTs……………………………………………..…203

5.3.7 Phase transformation…………………………………………………211

5.3.8 Aligned arrays of TiO2 NTs …………………………………………214

5.4 Combined sites for catalyzing SiO2 and TiO2 deposition…………….……218

5.4.1 Synthesis of the biomimetic catalyst…………………………………219

5.4.2 Synthesis of SiO2 catalyzed by the combined sites……………….…220

5.4.3 Synthesis of TiO2 catalyzed by the combined sites……………….…221

5.5. Conclusion…………………………………………………………………223

5.6 References…………………………………………………………….……225

Chapter 6 Mineralization of peptide self-assembled hydrogels…………….…227

6.1 Introduction……………………………………………………………..…227

6.2 Enzymatic self-assembly of Fmoc-Y and Fmoc-FY hydrogels……………227

6.2.1 Fmoc-Y hydrogel…………………………………………………….227

6.2.2 Fmoc-FY hydrogel………………………………………………...…229

6.3 Silicification of hydrogel nanostructures…………………………………..231

6.3.1 Silicification of Fmoc-Y gel…………………………………………231

6.3.1.1 Silicification via vortexing TEOS in the diluted hydrogels

(Method 1)………………………………………………………….……231

6.3.1.2 Silicification via depositing TEOS/H2O mixture on hydrogels

(Method 2)………………………………………………...…….…….…234

6.4 Conclusion………………………………………………………………….238

6.5 References …………………………………………………………………239

Chapter 7 Graphene-Inorganic hybrids……………………………………...…240

7.1 GO-Inorganic nanohybrids…………………………………………………240

7.1.1 Characterization of GO dispersion…………………………….……240

7.1.2 Preparation of GO-TiO2 nanohybrids ………………………………242

7.1.3 Preparation of GO-SiO2 nanohybrids…………………………….…247

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7.2 bwGO-Inorganic nanohybrids………………………………………...……251

7.2.1 bwGO dispersion……………………………………………………251

7.2.2 bwGO-TiO2 nanohybrids……………………………………...……253

7.2.2.1 Reaction in aqueous solution……………………………….……253

7.2.2.2 Reaction in EtOH…………………………………………..……255

7.3 Exfoliated graphene-Inorganic nanohybrids…………………………….…260

7.3.1 Effect of sonication time and centrifuge speed on the concentration of

the graphene dispersion ……………………………………………260

7.3.2 Evidence for exfoliation to graphene ………………………...……262

7.3.2.1 Raman characterization of the exfoliated samples…………..…262

7.3.2.2 TEM characterization of the exfoliated samples ………………273

7.3.2.3 AFM characterization of the exfoliated samples…………….…278

7.3.3 Preparation of exfoliated graphene (EG)-TiO2 nanohybrids………284

7.3.3.1 Preparation of EG-TiO2 hybrids in aqueous solution…………..284

7.3.3.2 Preparation of EG-TiO2 nanohybrids in EtOH…………………287

7.4 Conclusions………………………………………………………………...288

7.5 References……………………………………………………………….…290

Chapter 8 General conclusions and future work ………………………………293

8.1 General conclusions …………………………………………………….…293

8.2 Recommendation for future work …………………………………………297

8.3 References…………………………………………………………….....…297

Total Word Count: 65700

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List of tables

Table 3.1 Calculated molar absorptivity ε for all the modifiers studied………134

Table 3.2 Conditions used for the preparation of graphene dispersions……….147

Table 4.1 Initial adsorption rate of the Fmoc-AAs on CNT arrays……………170

Table 4.2 Calculated adsorption capacity (k) and intensity (n) for Fmoc-AAs

adsorbed on CNT arrays. Note that the units for k depend on the value of n. The

quality of fit, R2, was also given for each Fmoc-AA…………………………..173

Table 5.1 Measured SiO2 coating thickness based on TEM images…………...187

Table 5.2 Correlation of the adsorption equilibrium of the Fmoc-AAs on CNT

mats with the morphology of the hybrids ……………………………………..189

Table 5.3 Measured thickness of the TiO2 coating based on the TEM

observation……………………………………………………………………..201

Table 5.4 Measured inner diameter of the synthesized TiO2 NTs……………...205

Table 5.5 Measured wall thickness of the synthesized TiO2 NTs……………...205

Table 5.6 Measured inner diameter and wall thickness of the resultant TiO2

NTs …………………………………………………………………….………218

Table 7.1 Measured concentrations of graphene dispersions produced with

various sonication time and centrifuge speed …………………………………262

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List of figures

Figure 2.1 Schematic representation of sol-gel process of synthesis of

nanomaterials 7…………………………………………………………………..41

Figure 2.2 The structures of (a) SWNTs and (b) MWNTs 15……………………44

Figure 2.3 Schematic representation of a 2D graphene sheet with the lattice

vectors a1 and a2 and the roll-up vector Ch=na1+ma2. 18 …….………………….45

Figure 2.4 1,3-dipolar cycloaddition of an aminoethylene glycol linker to the

external surface of CNTs and the derivatization with N-protected glycine was

then obtained via amidation reaction. 85................................................................49

Figure 2.5 Fabrication of a glucose biosensor based on CNT nanoelectrode

ensembles 89……………………………………………………………………..50

Figure 2.6 Amine groups on a protein react with the anchored succinimidyl ester

to form amide bonds for protein immobilization.62…………………………….52

Figure 2.7(a) TEM micrographs of MWNTs dispersed with Fmoc-W (trp).

Arrows indicate the edge of the lattice structure upon which Fmoc-W aggregates

are apparent; (b) Optimized structures of (i) Fmoc-G (gly) and (ii) Fmoc-W

bound to [6,6] SWNTs with close-up images that highlight the orientation and

arrangement of Fmoc and the aromatic W ring 47……………………………….53

Figure 2.8 (i) SEM images of nano-1/SWNT fibres formed from a 100 μM

peptide/nanotube dispersion upon addition of no salt (A), 40 mM NaCl (B), and

120 mM NaCl (C). (ii) (A) SEM image of fibres formed from the addition of

0.0015% (by volume) DMF to a nano-1/SWNT dispersion. (B) Low-resolution

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TEM image of the same fibres observed in i(A). The small dark spheres are Fe

catalyst particles from the HiPco SWNT synthesis. (C) High-resolution TEM

image of the same fibres showing alignment of nanotubes. The large dark areas

are Fe particles 49. ………………………………………………………………54

Figure 2.9 Proposed mechanism of nanotube isolation from bundle 120…..……56

Figure 2.10 Schematic representations of the mechanisms by which surfactants

help disperse SWNTs. 101…………………………………………………..……57

Figure 2.11. Scheme for the preparation of CNT–silica nanohybrids.196……..…62

Figure 2.12 Scheme of the reaction between MWCNT-OH and AEAPS for the

following synthesis of silica coated MWCNTs. 197……………………..……….63

Figure 2.13 Mechanism of photocatalysis on the surface of TiO2 in presence of

UV radiation. 216…………………………………………………………………64

Figure 2.14 Schematic representation of a dye-sensitized solar cell based on

particulate TiO2. 217 ……………………………………………………….……..65

Figure 2.15 Schematic representation of the electron path through a (a) percolated

and (b) oriented nanostructure. 220 …………………………….………………..66

Figure 2.16 Electron transport across nanostructured semiconductor films: (A) in

the absence and (B) in the presence of CNTs support. 222……………..………..67

Figure 2.17 Left: Scheme of the beneficial role of benzyl alcohol in the in situ

coating of pristine CNTs with TiO2. One possible conformation of two BA molecules

on the CNT surface is shown in Scheme. Right: SEM images of TiO2 on CNTs after

conversion from anatase to rutile: A) no BA and B) with BA.172…………………..68

11

Figure 2.18 (a) Primary structures of the K2 and (QL)6 series of peptides showing

the comparative domain size. (b) Proposed model of nanofibre self-assembly

indicating hydrophobic packing region, axis of hydrogen bonding, and repulsive

positive charges. 271…………………………………………………………..…74

Figure 2.19 Computer modelling of the designed self-assembling fibre 274….…75

Figure 2.20 Potential pathway of V6D peptide nanotube formation.279…………77

Figure 2.21 (A) chemical structure of a PA which includes three distinct regions:

a hydrophobic alkyl tail, a glycine containing region, and a charged head group.

(B) Three-dimensional representation of the regions within the PA nanofibre.

Region (a) is the hydrophobic core composed of aliphatic tails. Region (b) is the

critical β-sheet hydrogen bonding portion of the peptide. Region (c) is the

peripheral peptide region which is not constrained to a particular hydrogen

bonding motif and forms the interface with the environment. 282………………79

Figure 2.22 Some of the possible modes of π-π interactions that contribute to the

emissions in the gel phase. 289………………………………………………..…80

Figure 2.23 (A) A model structure was created of Fmoc-FF peptides arranged into an

anti-parallel β-sheet pattern (i) which then come together through π–π interactions

between the Fmoc groups (in orange) (ii) like a zipper to create a cylindrical structure

(iii & iv) (B) TEM image of the Fmoc-FF hydrogels composed of flat ribbons made

up of side-by-side packing of the fibrils. 292…………………………………….....81

Figure 2.24 (A) Suspension of Fmoc-Leu2-OMe and inversion of glass vial

demonstrates self-supporting gel formation of Fmoc-Leu2 after ester hydrolysis

using subtilisin (Entry 1). (B) Proposed mechanism of Fmoc-peptide ester

hydrolysis that self-assembles to form higher-order aggregates through π–π

interlocked β-sheets. 305…………………………………………………………83

12

Figure 2.25 Solutions of Fmoc-Thr-OH and Leu-OMe. The inversion of the glass

vial demonstrates self-supporting gel formation of Fmoc-Thr-Leu-OMe via

reversed hydrolysis by thermolysin (entry 6). 305……………………………….83

Figure 2.26 (i) Chemical structure of Nap-FFGEY. (ii) Reversible modification of

the peptide gelator by a phosphatase/kinase reaction. (iii) Optical images of (A)

gel formed initially (B) the solution obtained after adding a kinase to A (C) gel

restored after adding a phosphatase to B. 306…………………………………….84

Figure 2.27 Proposed mechanism of silicon ethoxide condensation catalyzed by

silicatein α. 316…………………………………………………………………...87

Figure 2.28 Proposed condensation reaction between silicic acid and serine on the

protein template of the silicalemma. Water by-product may be eliminated or

structurally incorporated into the forming frustule through hydrogen bonding with

the oxygens of silica. 318………………………………………………………..88

Figure 2.29 Schematic of the interaction between two GNPs (B,C) capped with

imidazole and hydroxyl functionalities (A). (D) TEM image of silica product with

entrapped GNPs. Selected area electron diffraction (inset) indicating amorphous

nature of silica. 337……………………………………………………………….91

Figure 2.30 Mother of all graphitic forms. Graphene is a 2D building material for

carbon materials of other dimensionalities.338…………………….…….………92

Figure 2.31 Preparation of graphene by chemical reduction of GO synthesized by

Hummers’ method. …………………………………………………..…………97

Figure 2.32 Schematic model of a GO sheet, with -COOH hanging on the edge

and -O- and –OH decorate the basal plane. 388………………………………….98

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Figure 2.33 Schematic representation of as-produced GO: large oxidatively

functionalized graphene-like sheets with surface-bound debris. Note that the

graphene-like sheets extend further than depicted. 394……………..…………….98

Figure 2.34 TiO2-graphene composite and its response under UV-excitation.427…104

Figure 3.1 Schematic diagram showing the set-up for the CVD synthesis of

aligned CNT arrays………………………………………………………….…128

Figure 3.2 (a) Molecular structures of the modifiers studied. (b) Scheme

illustrating the UV-Vis measurement of the adsorption of the surfactant on (c)

aligned CNT arrays (side-view) and (d) randomly aligned CNT networks……130

Figure 3.3 Calibration curves of all the modifiers studied.……………………131

Figure 3.4 Schematic illustration of the competitive binding from the library

solution of Fmoc-AAs on graphite……………………………….…………….137

Figure 3.5 Molecular structure of THEOS……………………..………………145

Figure 3.6 Schematic diagram of a TEM. 11 ………………………..…………151

Figure 3.7 Ray path in a TEM operating in (a) image mode (b) diffraction mode.

12………………………………………………………………………………..152

Figure 3.8 Schematic representation of reversed-phase HPLC. The most

hydrophilic components (orange) elute from the column first, followed by the less

hydrophilic components (green), and finally the most hydrophobic components

(blue). 13……………………………………………………………….………..154

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Figure 3.9 Schematic of a liquid drop on a solid surface, where the

solid–vapor interfacial energy is denoted by γsv, the solid–liquid interfacial

energy is denoted by γsl, and the liquid–vapor interfacial energy is denoted by

γlv. 14 ………………………………………………………………….………..155

Figure 3.10 Sessile drop method for determining the contact angle. The fitted

contour is shown in green. 15………………………………..…………………156

Figure 3.11 (a) Typical Raman spectra for bulk graphite and monolayer graphene

obtained using a 514 nm laser. (b) Comparison of the D band at 514 nm at the

edge of bulk graphite and monolayer graphene. The fit of D1 and D2 components

of the D band of bulk graphite is shown. 18……………………………………157

Figure 3.12 Measured 2D band for (a) monolayer, (b) bilayer, (c) trilayer, (d)

four-layer and (e) HOPG using a 514 nm laser. 20……………………………..158

Figure 3.13 Schematic diagram of the beam deflection system in an atomic force

microscope, using laser and photodetector to measure the beam position. 25 …160

Figure 4.1 SEM images of CNT arrays grown at 760 ºC from a 5wt% ferrocene in

toluene solution on SiO2 substrate for 1h. (a) Cross-sectional image of the aligned

CNT arrays. (b) Close-up view of the CNTs from the arrays. (c) TEM image of

the pristine CNTs with dark particles presented both in the hollow cavity and the

walls of CNTs (indicated by arrows). Scale bar, 0.2 μm. (d) HRTEM image

showing the multilayered structure of a synthesized CNT with the lattice fringes

clearly visible. Scale bar, 5 nm. (e) The corresponding SAED pattern was indexed

to the (002), (100) and (004) planes of MWNTs……..…………..…………….165

Figure 4.2 Adsorption profiles of (a) Fmoc-Trp (c) Fmoc-Phe (e) Fmoc-Tyr (g)

Fmoc-His (i) Fmoc-Gly and (l) BA on aligned CNT arrays. (b,d,f,h,j)

Determination of the initial adsorption rate of the corresponding modifiers on the

15

arrays. (k) Histogram showing the equilibrium loadings of the Fmoc-AAs on the

arrays.…………………………………………...……………………...………168

Figure 4.3 Adsorption profile of Fmoc-Trp on randomly aligned CNT

networks………………………………………………………………………..170

Figure 4.4 Desorption profiles of (a) Fmoc-Trp and (b) Fmoc-Phe from CNT

arrays in water…………………………………………………………………171

Figure 4.5 Plot of ln Q vs. ln C for the adsorption of Fmoc-Trp (red circles) and

Fmoc-Gly (blue triangles) on the arrays……………………………………….173

Figure 4.6 HPLC chromatogram of 0.4 mM of (a) Fmoc-Phe (b) Fmoc-Trp (c)

Fmoc-Tyr (d) Fmoc-Gly and (e) Fmoc-His. (f) The mixture of the 5 Fmoc-AAs

with the same volume ratio……………………………………………………175

Figure 4.7 (a) HPLC traces of the mixture consisting of the five Fmoc-AAs at 0 h

(upper) and after 173 h of competitive binding (lower). (b) Comparison of the

equilibrium loadings of the five Fmoc-AAs on graphite in individual adsorption

and competitive binding experiments………………………………..…………175

Figure 4.8 Displacement of Fmoc-Gly by Fmoc-Trp on HOPG surface………178

Figure 5.1 SEM images of (a) the product obtained from the control experiment

in which pristine CNTs were used as templates. (b) Silica coated Fmoc-Trp and (c)

Fmoc-His functionalized CNTs. (d) A mixture of partially coated and uncoated

CNTs in the presence of Fmoc-Tyr after reaction for 21 days. (e) EDX spectrum

of the product shown in (c). Note that the aluminum and some of the oxygen were

from the sample stub………………………………....………………………...183

Figure 5.2 TEM images of (a) pristine CNTs co-existed with isolated SiO2

16

particles. Note. The image was over-focused as it was taken during early stage of

the PhD. Silica coated Fmoc-Trp functionalized CNTs after reaction for (b) 3

days and (c) 21 days. Silica coated Fmoc-His functionalized CNTs after reaction

for (d) 3 days and (e) 21 days. Partially coated Fmoc-Tyr functionalized CNTs

after reaction for (f) 3 days and (g) 21 days. Scale bar, (a) 100nm, (b) 20nm,

(c)-(g) 50nm……….……..……………………………………………….……186

Figure 5.3 (a) Line profile taken perpendicular to the tube axis direction. Inset:

Dark field STEM image of the hybrid NT. The direction of the scan was marked

by the arrow. The analysis was conducted with the help of Xiaofeng Zhao. (b)

Cross sectional view of a SiO2 coated CNT. The interaction of electron beam

with the edge and the centre of the hybrid tube was indicated by the red and

yellow line respectively. Blue colour: silica coating…………….…………..…188

Figure 5.4 Proposed catalytic mechanisms for silica templating………………190

Figure 5.5 SEM images of silica coated Fmoc-His functionalized CNTs obtained

after a growth time of (a) 3 days (b) 7 days and (c) 21 days. (d) Plot of the

diameter of the hybrid NT against the growth time. The average value was

calculated based on 50 separate measurements..………………..…………..…192

Figure 5.6 TEM images of silica coated Fmoc-Trp functionalized CNTs (a) before

and (b) after annealing at 200°C, and silica coated Fmoc-His functionalized CNTs

(c) before and (d) after annealing under the same

condition……………………………………………….………………………193

Figure 5.7 SEM images of (a) the product obtained using pristine CNTs as

templates. TiO2 coated CNTs in the presence of (b) Fmoc-Trp (c) Fmoc-His (d)

Fmoc-Tyr and (e) BA. (f-h) EDX spectra measured for the hybrids shown in (b-d).

Note the Al signal was originated from SEM stub, and Pt signal was originated

from the conductive coating on the SEM sample to reduce charging effect. The

17

considerably stronger C signal in (h) was due to the application of a thin layer of

carbon on the SEM sample as the conductive coating………….……..………195

Figure 5.8 TEM images of (a) the product obtained using pristine CNTs as

templates. TiO2 coated CNTs in the presence of Fmoc-Trp with the CNT

concentration of (b) 30 wt% and (c) 12 wt%. TiO2 coated CNTs in the presence

of Fmoc-His with the CNT concentration of (d) 30 wt% and (e) 12 wt%. A

cluster of TiO2 nanoparticles were deposited on the smooth surface of the TiO2

coating in (e). TiO2 coated CNTs in the presence of Fmoc-Tyr with the CNT

concentration of (f) 30 wt% and (g) 12 wt%. TiO2 coated CNTs in the presence of

BA with the CNT concentration of (h) 30 wt% and (i) 12 wt%. The arrows

indicated the uncoated part of CNTs. Note. This was different from the cracks

resulting from the drying effect. (j) SAED pattern taken from the sample shown

in (f). (k) XRD pattern of the as-produced CNT-TiO2 nanohybrids. C: CNT. For

(c), (e), (g) and (i), scale bar = 200 nm. For (a), (b), (d), (f) and (h), scale bar =

100 nm.….……………………………………………...………………………199

Figure 5.9 SEM images of the structures produced with the addition of (a)

undiluted and (b) diluted Fmoc-His solutions (by a factor of 10)…………….202

Figure 5.10 SEM images of TiO2 coating growing on Fmoc-Trp functionalized

CNTs at different reaction times of (a) 10 min (b) 1 h and (c) 6.5 h. (d) Plot of the


calculated based on 50 separate measurements……..……….…………………203

Figure 5.11 SEM images of TiO2 nanotubes produced from (a) TiO2 coated

Fmoc-His functionalized CNTs (30 wt%) and (b) TiO2 coated Fmoc-Tyr

functionalized CNTs (12 wt%). (c) EDX spectrum of the hybrid after calcination

at 550 ºC. Note. Pt signal was originated from the conductive coating on the SEM

sample. Scale bar, (a) 500nm, (b) 1μm…...……………………………………204

18

Figure 5.12 TEM images of the calcined hybrids. (a) In the presence of Fmoc-Trp

and 30wt% of CNTs. (b) In the presence of Fmoc-Trp and 12wt% of CNTs. (c) In

the presence of Fmoc-His and 30wt% of CNTs. (d) In the presence of Fmoc-His

and 12wt% of CNTs. (e) In the presence of Fmoc-Tyr and 30wt% of CNTs. (f) In

the presence of Fmoc-Tyr and 12wt% of CNTs. (g) In the presence of BA and

30wt% of CNTs. (h) In the presence of BA and 12wt% of CNTs. (i-l) SAED

patterns taken from the samples shown in (b-d) and (f) respectively (upper half)

which confirmed the polycrystalline anatase phase of the NTs by showing

excellent agreement with those simulated from JCPDS 21-1272 (lower half). The

SAED patterns were indexed to the (101), (004), (200) and (211) planes of

anatase phase. (m) XRD pattern taken from the sample shown in (d). A: anatase.

For (a), (e) and (g), scale bar = 100 nm and for (b), (c), (d), (f) and (h), scale bar =

200 nm……...………………………………………………………………….208

Figure 5.13 HRTEM image of a synthesized TiO2 NT showing the lattice spacing

of 0.35 nm, corresponding to the (101) crystal planes of anatase. Scale bar,

10nm……………………………………………………………………………210

Figure 5.14 (a) TEM images of the hybrids after heat treatment in Ar at 900 ºC

followed by in air at 550 ºC with the ramp rate of 20 ºC/min. Scale bar, 20 nm. (b)

XRD pattern taken from the sample shown in (a). (c) TEM image of the hybrids

after heat treatment in Ar at 800 ºC followed by in air at 550 ºC with the ramp

rate of 20 ºC/min. Scale bar, 100 nm. (d) SAED pattern (upper half) taken from

the sample shown in (c). The pattern was indexed to the (101), (004), (200) and

(211) planes of anatase phase. (e) TEM image of the hybrids after heat treatment

in air at 400 ºC followed by in Ar at 800 ºC with a ramp rate of 20 ºC/min. Scale

bar, 100 nm. (f) XRD pattern taken from the sample shown in (e). A: anatase, R:

rutile, C: CNT. (g) TEM image of the hybrids after heat treatment in air at 400 ºC

followed by in Ar at 800 ºC with a ramp rate of 1 ºC/min. Scale bar, 200 nm. (h)

SAED pattern (upper half) taken from the sample shown in (g). The SAED

19

pattern was indexed to the (110), (111), (210), (211) and (220) planes of rutile

phase. ……………………………………………….………………..…….…..213


where as-produced CNT mat was used as templates. TiO2 NT arrays produced in

the presence of (b) Fmoc-Trp (c) Fmoc-His (d) Fmoc-Tyr and (e) BA……….216

Figure 5.16 TEM images of (a) the product obtained from the control experiment.

TiO2 NTs produced in the presence of (b) Fmoc-Trp (c) Fmoc-His and (d)

Fmoc-Tyr. (e) Collapsed NT structures obtained in the presence of BA. The red

arrow in (b) and (c) indicated the open ends of the TiO2 NTs. Note. CNT

templates were not completely removed after calcination as indicated by the black

arrows in (c). (f) XRD pattern taken from the sample shown in (b). Scale bar, (a-e)

200 nm………………………………………………………………………….217

Figure 5.17 SEM images showing (a) bundled fibers and (b) spherical aggregates

formed in the combined solutions. (c,d) Magnified images of the aggregates

shown in (a) and (b) respectively. (e) Fmoc-His f-CNTs and (f) Fmoc-Tyr

f-CNTs…………………………………………………………………………219

Figure 5.18 (a,b) SEM image of silica coated combined catalyst after heat

treatment. (c) EDX spectrum of the sample shown in (a)……………...………221

Figure 5.19 SEM images of (a) TiO2 nanorods coated CNT bundles (b) TiO2

nanorods coated individual CNTs (c) TiO2 nanorods coated CNT bundles after

heat treatment and (d) TiO2 particles formed on Si wafer. (e) and (f) EDX

spectrum of the sample shown in (a) and (c) respectively…………….……….222

Figure 6.1 (a) Schematic representation of the enzymatic dephosphorylation of

Fmoc-Y(p)-OH to Fmoc-Y. The corresponding optical images for Fmoc-Y(p)-OH

precursor solution before enzyme addition and the self-supporting hydrogels

20

formed were also shown. (b) Negatively stained TEM image of the diluted

Fmoc-Y hydrogel. (c,d) Negatively stained TEM image of the undiluted

hydrogel.. …………………………………………………………..…………..228

Figure 6.2 (a) AP catalyzed dephosphorylation reaction of Fmoc-FpY and a

schematic representation of the supramolecular transition from micelles to fibres2.

(b) Negative stained TEM image showing the Fmoc-FY self-assembled

nanofibrils. Scale bar, 100 nm. (c) HPLC trace of the conversion of Fmoc-FpY to

Fmoc- FY as a function of time. The gelation point is marked with an arrow. (d)

Fluorescence emission spectra of the solution of Fmoc-FpY and the hydrogel of

Fmoc-FY………………………………………………………………………230

Figure 6.3 TEM images of silica coating on Fmoc-Y self-assembled

nanostructures after reaction for (a) 1 h, (b) 2 h, (c,d) 4 h and (e) 5 h. Scale bar,

100 nm. (f) EDX spectrum of the mineralized peptide nanofibrils. (f) EDX

spectrum of the silicified fibrils…………………………………………...……232

Figure 6.4 Silicification process of Fmoc-Y hydrogel…………………………234

Figure 6.5 SEM analysis on (a) the upper aqueous phase and (c) the lower

hydrogel phase. (b) EDX spectrum of (a)……………..……………………….236

Figure 6.6 Unstained TEM images of (a) the network of silicified hydrogel

nanofibrils that were derived from the resulting clear gel. Scale bar, 100 nm. (b)

Fmoc-Y self-assembled hydrogel. Scale bar, 200 nm …………………………238

Figure 7.1 (a) SEM image of aggregated GO sheets. (b) TEM image of single

layer GO sheet with folds present at both sides (indicated by arrows). Scale bar,

100 nm. (c) Corresponding SAED pattern taken from the region marked by the

dashed box in (b). The pattern was labeled with Miller-Bravais indices. (d)

Intensity profile plot along the line between the arrows shown in (c). (e) Lower

21

magnification TEM image of GO sheets with the folds indicated by arrows. Scale

bar, 200 nm. (f) Corresponding SAED pattern taken from the region marked by

the dashed box in (e) showing three superimposed hexagonal patterns indicated

by yellow, red and blue colors………………………………………………….241

Figure 7.2 (a) TEM image of GO-TiO2 nanohybrids produced with lower TBOT

concentration for 4 h. Inset corresponds to the SAED pattern taken from the

region marked by the red dashed box. (b) A magnified image of the region shown

in the orange dashed box in (a). (c) EDX spectrum of (a). Note that Cu signal is

originated from the TEM grid. (d) TEM image of the hybrids produced with

lower TBOT concentration for 7 d. (e) TEM image of the hybrids produced with

higher TBOT concentration for 4 h. (f) Corresponding SAED pattern taken from

the region marked by the dashed box in (e) and the diffraction spots are labeled

using Miller-Bravais indices. (g) Intensity profile plot along the line between the

arrows shown in (f).…………………………………………………………….244

Figure 7.3 TEM images of the thermally treated nanohybrids obtained from (a)

the reaction with lower TBOT concentration for 4h and (b) the reaction with

higher TBOT concentration for 4h. The inset in (a) and (b) showed the

corresponding SAED patterns which were indexed to (c,e) GO (labeled using

Miller (hkl) indices) and (d,f) anatase TiO2 respectively. Note that the upper half

in (c)-(f) showed the experimental data while the lower half in (c) and (e) showed

the diffraction pattern of GO, and that in (d) and (f) showed the simulated

diffractions for anatase according to JCPDS 21-1272…………………………246

Figure 7.4 SEM images of (a) highly aggregated GO sheets. (b) GO-SiO2

nanohybrids with layered structure (indicated by arrows along the edges). (c)

Higher magnification image showing the partial separation of two hybrid sheets.

(d) EDX spectrum of the sample shown in (b)…………………………………248

22

Figure 7.5 (a) Low magnification TEM image of silica coated GO sheets. (b) A

magnified TEM image showing the ripples present on the GO sheet (indicated by

the arrow). The SAED pattern taken from the region marked by the dashed box

was labeled using Miller-Bravais indices………………………………………249

Figure 7.6 (a) TEM image of porous silica sheets obtained from the calcination of

GO-SiO2 hybrids. (b) Corresponding SAED pattern taken from the sample shown

in (a)……………………………………………………………………………250

Figure 7.7 Photographs of the aqueous dispersion of bwGO (a) in the absence and

(b) in the presence of Fmoc-Trp. The dispersions were allowed to stand for 35

days…………………………………………………………………………….252

Figure 7.8 (a) TEM image of bwGO sheets deposited from the dispersion in

Fmoc-Trp solution. (b) The corresponding SAED pattern taken from the sample

shown in (a). The pattern was labeled using Miller (hkl) indices……………...253

Figure 7.9 TEM image of bwGO-TiO2 nanohybrids prepared in aqueous solution.

The arrows indicate the wrinkles present in bwGO sheets. (b) The corresponding

SAED pattern taken from the sample shown in (a)……………………………254

Figure 7.10 Raman spectra for (a) bwGO deposited from the dispersion in

Fmoc-Trp solution (b) anatase TiO2 and (c) annealed bwGO-TiO2 nanohybrids

prepared in aqueous solution. The spectra were taken using a 633 nm HeNe laser.

Note that the peak at around 520 cm-1 was attributed to the SiO2/Si

substrate…………………………………………………………………….…..255

Figure 7.11 (a) SEM image of bwGO-TiO2 nanohybrids prepared in EtOH. (b)

EDX spectrum. Pt signal is originated from Pt coating on the SEM sample to

reduce charging effect……….…………………………………………………256

23

Figure 7.12 (a) TEM image of bwGO-TiO2 nanohybrids prepared in EtOH with

the addition of H2O. (b) Corresponding SAED pattern taken from the region

marked by the dashed box in (a). (c) TEM image of bwGO-TiO2 nanohybrids

prepared in EtOH with the addition of Fmoc-Trp solution. (d) Corresponding

SAED pattern taken from the region marked by the dashed box in (c). The pattern

was labeled using Miller-Bravais indices. (e) Intensity profile plot along the line

between the arrows shown in (d)………………………………………………257

Figure 7.13 Schematic illustration of the synthesis of bwGO-TiO2 nanohybrids in

(a) EtOH and (b) aqueous solution…………………………………………..…258



prepared in EtOH. The spectra were taken using a 633 nm HeNe laser………260

Figure 7.15 Digital images of the graphene dispersions prepared under various

conditions………………………………………………………………………261

Figure 7.16 Raman spectra for (a) the starting graphite powder and the flakes

deposited from the dispersions prepared with (b) 1 h (c) 6 h and (d) 12 h of

sonication followed by centrifugation at 3000 rpm respectively. The spectra were

measured on SiO2/Si substrate and in all cases the excitation wavelength was

633nm. D, G, 2D and D’ bands are indicated in the Figure. All the spectra were

normalized to have the similar G band intensity and offset for

clarity…………………………………….…………………………….……….264


deposited from the dispersions prepared with centrifugation at (b) 500 rpm (c)

3000 rpm and (d) 6000 rpm following 6 h of sonication respectively. The spectra

were measured on SiO2/Si substrate and in all cases the excitation wavelength

was 633 nm. D, G, 2D and D’ bands are indicated in the Figure. All the spectra

24

were normalized to have the similar G band intensity and offset for

clarity………………………………………………………………………...…265

Figure 7.18 Histograms and normal distribution of the 2D band position for

varying sonication time and centrifuge speed…………………………………266

Figure 7.19 Mean 2D band position as a function of (a) sonication time and (b)

centrifuge speed. The data for the starting graphite powder was also shown for

comparison……………………………………………………………………267

Figure 7.20 Histograms and normal distribution of the 2D bandwidth for varying

sonication time and centrifuge speed…………………………………………267

Figure 7.21 Mean 2D bandwidth as a function of (a) sonication time and (b)

centrifuge speed………………………………………………………….……268

Figure 7.22 Histograms and normal distribution of I2D/IG ratio for varying


Figure 7.23 Mean I2D/IG ratio as a function of (a) sonication time and (b)

centrifuge speed………………………………………………………………270

Figure 7.24 Histograms and normal distribution of ID/IG ratio for varying


Figure 7.25 Mean ID/IG ratio as a function of (a) sonication time and (b)

centrifuge speed. The ratio for the starting graphite was also shown for

comparison…………………………………………………………………….272

Figure 7.26 Plot of ID/IG ratio against 2D band position for varying sonication

time and centrifuge speed. The data for the starting graphite was also shown for

25

comparison. The direction of the arrow corresponds to flakes of fewer layer and

smaller size……………………………………………………………………273

Figure 7.27 Representative TEM images of graphene flakes deposited from the

dispersions prepared with various sonication time and centrifugation speed….274

Figure 7.28 (a-e) Histograms and normal distribution of the flake area for varying

sonication time and centrifuge speed. (f) Mean flake area as a function of

sonication time. (g) Mean flake area as a function of centrifuge speed………277

Figure 7.29 AFM characterization of the exfoliated flakes……………………281

Figure 7.30 TEM images of the EG-TiO2 nanocomposites……………………285

Figure 7.31 TEM images of EG-TiO2 nanohybrids prepared in EtOH………..287

26

List of abbreviations

1D 1- dimensional

2D 2-dimensional

3D 3- dimensional

AFM Atomic force microscopy

Al Aluminum

Ala or A Alanine

ALD Atomic layer deposition

AP Alkaline phosphatase

APTES 3-aminopropyltriethoxyysilane

Asn or N Asparagine

Asp or D Aspartic acid

BA Benzyl alcohol

BSA Bovine serum albumin

bwGO Base-washed graphene oxide

C Carbon

CCG Chemically converted graphene

CD Circular dichroism

CMGs Chemically modified graphenes

CNT Carbon nanotube

Cu Copper

CVD Chemical vapour deposition

Cys or C Cysteine

D2O Deuterium oxide

dH2O Deionized H2O

DLS Dynamic light scattering

DMF Dimethylformamide

EDX Energy dispersive x-ray

EG Exfoliated graphene

EtOH Ethanol

FEGSEM Field emitter gun scanning electron microscope

FGSs Functionalized graphene sheets

Fmoc-AA N-(fluorenyl-9-methoxycarbonyl) terminated amino acid

Fmoc-FY Fmoc-Phenylalanine-Tyrosine

Fmoc-FpY Fmoc-Phenylalanine-Tyrosine (phosphate)

Fmoc-Y Fmoc-Tyrosine

Fmoc-Y(p)-OH Fmoc-Tyrosine (phosphate)-OH

FT-IR Fourier transform infrared spectroscopy

FWHM Full width at half maximum

Glu or E Glutamic acid

Gly or G Glycine

GNPs Golden nanoparticles

GO Graphene oxide

27

GS Graphene sheets

HA Hydroxyapatite

HeNe Helium–neon

HiPco High-pressure decomposition of carbon oxide

His or H Histidine

HOPG Highly orientated pyrolytic graphite

HPLC High performance liquid chromatography

HRTEM High-resolution transmission electron microscopy

H-bonding Hydrogen bonding

iTO in-plane transverse optical

Leu or L Leucine

LO longitudinal optical

Lys or K Lysince

MWNTs Multi-walled nanotubes

NaOH Sodium hydroxide

NMP N-Methyl-2-Pyrrolidone

NMR Nuclear magnetic resonance

NT Nanotube

O Oxygen

OD Oxidative debris

PA Peptide amphiphile

PECS Precision Etching Coating System

Phe or F Phenylalanine

Pt Platinum

QDs Quantum dots

RGO Reduced graphene oxide

rpm Revolutions per minute

SAED Selected area electron diffraction

SAF Self-assembling fibre

SDS Sodium dodecyl sulfate

SDBS Sodium dodecyl benzene sulfonate

SEM Scanning electron microscopy

Ser or S Serine

Si Silicon

SiO2 Silicon dioxide

STEM Scanning transmission electron microscopy

SWNTs Single-walled nanotubes

TBOT Tetrabutyl titanate

TEM Transmission electron microscopy

TEOS Tetraethyl orthosilicate

THEOS Tetrakis (2-hydroxyethyl) orthosilicate

Thr or T Threonine

Ti Titanium

TiO2 Titanium dioxide

Trp or W Tryptophan

Tyr or Y Tyrosine

https://www.google.co.uk/url?sa=t&rct=j&q=&esrc=s&source=web&cd=1&cad=rja&ved=0CC8QFjAA&url=http%3A%2F%2Fen.wikipedia.org%2Fwiki%2FScanning_transmission_electron_microscopy&ei=4dlEUraJH8ithQfrioH4BQ&usg=AFQjCNEBymaXBV7CmpOtUMcl4vRMcgrn4A

https://www.google.co.uk/url?sa=t&rct=j&q=&esrc=s&source=web&cd=1&cad=rja&ved=0CC0QFjAA&url=http%3A%2F%2Fen.wikipedia.org%2Fwiki%2FTetraethyl_orthosilicate&ei=P9pEUqi6DZCXhQf_jICIDQ&usg=AFQjCNEXXpcn2xoSRZns39TbXtXNKAy6fg

https://www.google.co.uk/url?sa=t&rct=j&q=&esrc=s&source=web&cd=1&cad=rja&ved=0CC0QFjAA&url=http%3A%2F%2Fen.wikipedia.org%2Fwiki%2FTitanium_dioxide&ei=X9pEUsywFc6UhQfV8YHACg&usg=AFQjCNFHDZ_wfBW2uwH3zY8mN1G61G_Bvg

28

UV-Vis Ultraviolet-visible light spectroscopy

Val or V Valine

wt% Weight%

XRD X-ray diffraction

29

List of Symbols

A Absorbance

a1, a2 Lattice vectors of graphene sheet

b Path length

C Equilibrium concentration of the solute in solution

c Concentration

Ch Chiral vector

d Crystal size

I Intensity

I2D Raman intensity for 2D band

IA Integrated intensity of anatase (101) peak

ID Raman intensity for D band

IG Raman intensity for G band

IR Integrated intensity of rutile (110) peak

K Shape factor

k Adsorption capacity constant

ki Initial adsorption rate

n Adsorption intensity constant

(n, m) Indices defining the nanotube structure

Q Amount of the solute adsorbed per unit weight of the

adsorbent

R2 Correlation coefficient

S Surface area

T Translation vector

t Time

V Volume

WR Percentage of rutile

β Full width at half maximum intensity

ε Molar absorptivity

λ Wavelength

θ Angle

30

Abstract

Silica and titania nanoparticles have been produced by using carbon nanotubes

(CNTs) and graphene as templates in a sol-gel reaction. A range of Fmoc

terminated amino-acids (Fmoc-AAs) were studied as surface modifiers to

encourage the templating on the nanocarbons. After annealing the deposited

structures, the carbon templates were either left in place to give hybrid structures

or oxidized to leave pure inorganic nanoparticles.

Absorption studies were initially conducted to identify Fmoc-AAs that would

bind well to the CNTs. Fmoc-Trp had the best affinity for CNTs out of the amino

acids studied. The fully reversible nature of the binding process was

demonstrated via the desorption of Fmoc-AAs from CNTs in water. The

equilibrium data were found to be well described by the Freundlich isotherm

model. The competitive binding from a library of Fmoc-AAs on graphite was

developed to efficiently identify the strongest binding candidate.

The synthesis of CNT-SiO2 and CNT-TiO2 nanohybrids were successfully

demonstrated. The morphology of the hybrids was found to be dependent on the

CNT:precursor and Fmoc-AA:CNT ratios. Fmoc-AAs were believed to play a

dual role: (1) electrostatically stabilizing the NT dispersion and (2) the

functionalities from the side chains of the amino acids providing binding sites for

SiO2 and TiO2 deposition. Uniform anatase nanotubes (NTs) were synthesized

after calcination of the CNT-TiO2 nanohybrids. Both the inner diameter and wall

thickness of the synthesized TiO2 NTs were controlled by the dimension of CNT

templates and the ratio of CNT:precursor. The transition from anatase to rutile

phase was found to be affected by heating temperature, pre-treatment and ramp

rate. A simple route towards the production of TiO2 NT arrays was also

demonstrated by using aligned CNT arrays as templates in the presence of the

Fmoc-AAs.

Graphene based nanohybrids were synthesized in the presence of graphene oxide

(GO), Fmoc-Trp stabilized base-washed graphene oxide (bwGO) and exfoliated

graphene via the sol-gel process. It was found that the morphology of the

products was highly dependent on the reaction media. Graphene dispersions were

prepared by direct exfoliation of graphite in Fmoc-Trp solution. Raman, TEM

and AFM analyses suggested the dispersion comprised of mainly few layer

graphene (<5 layers) with a broad size distribution and that the defects introduced

during sonication were predominately associated with the formation of new flake

edges due to sonication-induced cutting.

A preliminary study was conducted on the silicification of Fmoc-Y and Fmoc-FY

self-assembled hydrogels. The presence of a high density of –OH group on the

nanofibers’ surface was found to promote silica deposition.

31

Declaration

No portion of the work referred to in the thesis has been submitted in support of

an application for another degree or qualification of this or any other university or

other institute of learning.

32

Copyright Statement i. The author of this thesis (including any appendices and/or schedules to this

thesis) owns certain copyright or related rights in it (the “Copyright”) and s/he

has given The University of Manchester certain rights to use such Copyright,

including for administrative purposes.

ii. Copies of this thesis, either in full or in extracts and whether in hard or

electronic copy, may be made only in accordance with the Copyright, Designs

and Patents Act 1988 (as amended) and regulations issued under it or, where

appropriate, in accordance with licensing agreements which the University has

from time to time. This page must form part of any such copies made.

iii. The ownership of certain Copyright, patents, designs, trade marks and other

intellectual property (the “Intellectual Property”) and any reproductions of

copyright works in the thesis, for example graphs and tables (“Reproductions”),

which may be described in this thesis, may not be owned by the author and may

be owned by third parties. Such Intellectual Property and Reproductions cannot

and must not be made available for use without the prior written permission of

the owner(s) of the relevant Intellectual Property and/or Reproductions.

iv. Further information on the conditions under which disclosure, publication and

commercialisation of this thesis, the Copyright and any Intellectual Property

and/or Reproductions described in it may take place is available in the University

IP Policy (see http://documents.manchester.ac.uk/DocuInfo.aspx?DocID=487), in

any relevant Thesis restriction declarations deposited in the University Library,

The University Library’s regulations (see

http://www.manchester.ac.uk/library/aboutus/regulations) and in The University’s

policy on Presentation of Theses

33

Acknowledgements

The author would like to thank Prof. Ian Kinloch for his guidance and support

throughout. Further thanks go to Prof. Rein Ulijn who enabled the collaboration

and has offered guidance and encouragement at times of need.

Thanks also go to Chris, Polly, Alan, Gary, Andy, Xiaofeng for help with SEM,

UV-Vis, TEM, XRD, Raman, AFM and EDX linescan, Dr. Sarah Haigh for help

with TEM and discussion on diffraction pattern interpretation, Kate Thornton for

help with preparation of Fmoc-Tyr-OH hydrogel, Sangita Roy and Louise

Birchall for help with HPLC and fluorescence spectroscopy

Most of all, the author is indebted to her family for all their support and much

needed funding to complete the degree.

34

Chapter 1 Introduction

1.1 Overview Carbon nanotubes (CNTs) and more recently graphene have attracted

considerable interest owing to their unusual combination of electronic, thermal

and mechanical properties. Such remarkable properties have opened up a world of

possible applications, including photochemical, catalytic and electrochemical

technologies.

CNT-inorganic hybrid materials combine the physicochemical properties of

CNTs with the advantages of their inorganic components, leading to new

functionalities that do not exist in either building block 1,2. For example, the

hybrid materials exhibit a significant synergistic effect through size domain

effects and charge transfer processes across the CNT-inorganic interface.

Dielectric materials such as SiO2 and TiO2 are of particular interest amongst the

inorganic compounds. Owing to the high biocompatibility, hydrophilic nature and

easy surface functionalization of SiO2, CNT-SiO2 hybrids have found an

extensive range of applications, such as in biotechnology 3-5, nanoelectric devices

and reinforcement materials in composites 6-8. TiO2 has been extensively studied

as a highly active semiconductor photocatalyst material for applications in solar

energy conversion9, environmental purification10-20 and dye-sensitized solar

cells21-23. The combination of CNTs with the well-established photoactivity of

TiO2 has increased the charge-transfer efficiency, which further enhances the

photocatalytic activity 24-26.

Typically, inorganic compounds are coated onto CNTs using a sol-gel process

due to the mild reaction conditions (room temperate, near neutral pH etc.). The

morphology of the coating depends significantly on the surface chemistry of

CNTs, as the surface groups act as both catalysts and structural directors. Eder et

al. 27 employed benzyl alcohol as a surfactant to coat pristine CNTs with TiO2.

They assumed that the benzene ring of the surfactant adsorbed on CNT surface

35

via - stacking interactions, while the hydroxyl groups activated the hydrolysis

of the titanium precursor. Based upon this assumption, there should be a family of

surface modifiers which could enhance the inorganic templating process on the

nanotubes. For example, N-(fluorenyl-9-methoxycarbonyl) terminated amino

acids (Fmoc-AAs) are cheap and have previously been shown by the research

group 28 to bind well to CNTs and the amino acid group gives 26 different

functional motifs to explore for the templating reaction. The importance of the

amino acids is highlighted by studies on the biomineralization process of

silicateins in a marine sponge 29. The site-directed mutagenesis results have found

that both serine and histidine residues were required for the efficient catalysis of

the siloxane polymerization. Several synthetic counterparts have been developed.

For instance, peptide based self-assembled supramolecular structures have been

demonstrated to mimic the catalytic activity of silicateins for the templating of

silica 30, 31.

Graphene-TiO2 hybrid materials showed improved photocatalytic activity

compared with CNT-TiO2 attributed to the higher dye adsorption capacity and

enhanced charge separation and transportation properties. However, for such

applications to be achieved, suitable routes for graphene manufacturer have to be

developed. High-quality graphene has been produced by liquid-phase exfoliation

which includes the reduction of exfoliated GO 32,33 and sonication-assisted direct

exfoliation of graphite in solution34-39. Solvent exfoliation is particularly

attractive as it produces relatively defect-free graphene and can either be done in

a solvent such as NMP or in a surfactant solution. In particular for the latter, it

may be possible to select a surfactant that both enables exfoliation and can direct

templating.

1.2 Aims

Thus, this thesis initially aims to identify suitable Fmoc-AA surface modifiers for

CNTs and understand the adsorption process for these Fmoc-AA modifiers. The

36

Fmoc-AA coated CNTs will then be used as templates in the sol-gel deposition of

silica and titania, and their performance compared to that of the published benzyl

alcohol. The Fmoc-AA functionalized nanotubes also allow an attempt at

mimicking the catalytic active site of silicatein. The identified successful

Fmoc-AA will then be used as a surfactant to exfoliate graphene from graphite

and as a surface modifier in the production of inorganic-graphene hybrids using a

range of graphene materials.

More explicitly, the thesis aims to

(1) Study the non-covalent functionalization of CNTs through the adsorption of a

library of aromatic Fmoc-AAs on both aligned CNT arrays and randomly

aligned CNT networks.

(2) Synthesize silica and titania based nanohybrids via an in-situ sol-gel process

employing the Fmoc-AA functionalized CNTs as templates. Herein,

Fmoc-Trp, Fmoc-His and Fmoc-Tyr which render the templates’ surface

with the functionalities that have been reported to catalyze silica and titania

deposition were investigated as surface modifiers. The surface modifier is

expected to serve two purposes: helps to colloidally stabilize the CNT

dispersion as well as to promote the deposition of silica and titania on CNTs.

The role of the surface chemistry of CNTs in controlling the coating

morphology was also investigated.

(3) Prepare graphene dispersion by direct exfoliation of graphite in Fmoc-AA

solution and subsequently produce graphene–TiO2 nanohybrids employing the

exfoliated flakes as templates via sol-gel process. The degree of exfoliation

and quality of the exfoliated flakes was characterized by Raman spectroscopy,

TEM and AFM.

(4) Conduct a preliminary study on the silicification of Fmoc-Y and Fmoc-FY

37

self-assembled hydrogels. Both of the gels were prepared through an enzyme

catalyzed dephosphorylation. The presence of a high density of –OH group

on the nanofibers’ surface was expected to promote silica deposition.

1.3 References

1. Y. Zhang et al., Reinforcement of silica with single-walled carbon nanotubes

through covalent functionalization, J. Mater. Chem., 2006, 16, 4592.

2. M. Bottini et al., Non-destructive decoration of full-length multi-walled carbon

nanotubes with variable amounts of silica gel nanoparticles, Carbon, 2006, 44,

1301.

3. D. T. Mitchell et al., Smart Nanotubes for Bioseparations and Biocatalysis,

J. Am. Chem. Soc., 2002, 124, 11864.

4. S. Kapoor et al., Ultrasound-Triggered Controlled Drug Delivery and

Biosensing Using Silica Nanotubes, J. Phys. Chem. C, 2009,113, 7155.

5. C. C. Chen et al., Preparation of Fluorescent Silica Nanotubes and Their

Application in Gene Delivery, Adv.Mater., 2005, 17, 404.

6. M. Olek et al., Nanomechanical properties of silica-coated multiwall carbon

nanotubes-poly(methyl methacrylate) composites, Langmuir, 2005, 21, 3146.

7. J. Wang et al., Utilization of silane functionalized carbon nanotubes-silica

hybrids as novel reinforcing fillers for solution styrene butadiene rubber,

Polym. Compos., 2013, 34, 690.

8. W. Zhang et al., Polymer Nanocomposites Using Urchin-Shaped Carbon

Nanotube-Silica Hybrids as Reinforcing Fillers, Macromol. Rapid Commun.,

2004, 25, 1860.

9. S. Lee et al., Two-Step Sol-Gel Method-Based TiO2 Nanoparticles with

Uniform Morphology and Size for Efficient Photo-Energy Conversion Devices,

Chem. Mater., 2010, 22, 1958.

10. I. Sopyana et al., An efficient TiO2 thin-film photocatalyst: photocatalytic

properties in gas-phase acetaldehyde degradation, J. Photochem. Photobiol. A:

Chem., 1996, 98, 79.

11. Q. Li et al., Antimicrobial nanomaterials for water disinfection and microbial

control: potential applications and implications, Water Res., 2008, 42, 4591.

12. M. L. Sauer et al., Photocatalyzed Oxidation of Ethanol and Acetaldehyde in

Humidified Air, J. Catal., 1996, 158, 570.

13. T. N. Obee et al., TiO2 Photocatalysis for Indoor Air Applications: Effects of

Humidity and Trace Contaminant Levels on the Oxidation Rates of

Formaldehyde, Toluene, and 1,3-Butadiene, Environ. Sci. Technol., 1995, 29,

1223.

14. C.H. Aoa et al., Photodegradation of volatile organic compounds (VOCs) and

NO for indoor air purification using TiO2: promotion versus inhibition effect of

NO, Appl. Catal. B: Environ., 2003, 42, 119.

15. L. Petrov et al., Photocatalytic properties of modified TiO2 coatings for

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purification of waste water and air, J. Environ. Prot. Ecol., 2007, 8, 881.

17. I. Salema et al., Recent Studies on the Catalytic Activity of Titanium,

Zirconium, and Hafnium Oxides, Catal. Rev. Sci. Eng., 2003, 45, 205.

18. Z. Liu et al., Efficient Photocatalytic Degradation of Gaseous Acetaldehyde

by Highly Ordered TiO2 Nanotube Arrays, Environ. Sci. Technol., 2008, 42, 8547.

19. O. Carp et al., Photoinduced reactivity of titanium dioxide, Prog. Solid State

Chem., 2004, 32, 33.

20. T. Hirakawa et al., Mechanism of Photocatalytic Production of Active

Oxygens on Highly Crystalline TiO2 Particles by Means of Chemiluminescent

Probing and ESR Spectroscopy, J. Phys. Chem. B, 2001, 105, 6993.

21. M. Grätzel et al., A low-cost, high-efficiency solar cell based on

dye-sensitized colloidal TiO2 films, Nature, 1991, 353, 737.

22. M. Grätzel et al., Photoelectrochemical cells, Nature, 2001, 414, 338.

23. M. Adachi et al., Formation of Titania Nanotubes and Applications for

Dye-Sensitized Solar Cells, Journal of the Electrochemical Society, 2003, 150,

G488.

24. Y. Yu et al., Enhancement of photocatalytic activity of mesoporous TiO2 by

using carbon nanotubes, Appl. Catal. A: Gen., 2005, 289, 186.

25. W. Wang et al., Visible light photodegradation of phenol on

MWNT-TiO2 composite catalysts prepared by a modified sol–gel method, J. Mol.

Catal. A: Chem., 2005, 235, 194.

26. W. Wang et al., Photocatalytic degradation of phenol on MWNT and titania

composite catalysts prepared by a modified sol–gel method, Appl. Catal., B, 2005,

56, 305.

27. D. Eder et al., Carbon–Inorganic Hybrid Materials, Adv.Mater., 2008, 20,

1787.

28. B. G. Cousins et al., Enzyme-Activated Surfactants for Dispersion of Carbon

Nanotubes, Small, 2009, 5, 587.

29. Y. Zhou et al., Efficient catalysis of polysiloxane synthesis by silicatein a

requires specific hydroxyl and imidazole functionalities, Angew.Chem.Int.Ed.,

1999, 38, 780.

30. J. N. Cha et al., Biomimetic synthesis of ordered silica structures mediated by

block copolypeptides, Nature, 2000, 403, 289.

31. V. M. Yuwono et al., Peptide Amphiphile Nanofibers Template and Catalyze

Silica Nanotube Formation, Lagmuir, 2007, 23, 5033.

32. S. Stankovich et al., Synthesis of graphene-based nanosheets via chemical

reduction of exfoliated graphite oxide, Carbon, 2007, 45, 1558.

33. S. Dubin et al., A One-Step, Solvothermal Reduction Method for Producing

Reduced Graphene Oxide Dispersions in Organic Solvents, ACS Nano, 2010, 4,

3845.

34. J. N. Coleman et al., Liquid-Phase Exfoliation of Nanotubes and Graphene,

Adv. Funct.Mater., 2009, 19, 3680.

35. S. De et al., Flexible, transparent, conducting films of randomly stacked

graphene from surfactant-stabilized, oxide-free graphene dispersions, Small, 2010,

6, 458.

39

36. J. N. Coleman et al., High-yield production of graphene by liquid-phase

exfoliation of graphite, Nature Nanotechnol., 2008, 3, 563.

37. U. Khan et al., High-concentration solvent exfoliation of graphene, Small,

2010, 6, 864.

38. M. Lotya et al., Liquid Phase Production of Graphene by Exfoliation of

Graphite in Surfactant/Water Solutions, J. Am. Chem. Soc., 2009, 131, 3611.

39. M. Lotya et al., High-concentration, surfactant-stabilized graphene

dispersions, ACS Nano, 2010, 4, 3155.

40

Chapter 2 Literature review

2.1 Sol-gel chemistry

The sol–gel process is the most popular technique for the production of glasses

and ceramic materials due to its low reaction temperatures compared to melting

glass or firing ceramics. Also, as a wet-chemical technique, it has many

advantages over other conventional "powder" routes, including 1: (1) The

rheological properties of sols and gels allow the production of various forms of

products including ultrafine powders, thin films, fibers and monoliths depending

on the processing conditions 2-4. (2) Easy deposition of good quality coatings onto

a variety of substrates. (3) Better control over the whole process and the synthesis

of "tailor-made" materials. (4) Production of high-purity materials at mild

reaction conditions which are highly desired in some applications such as

bioencapsulation and sensors 5,6.

The sol-gel process involves the formation of a colloidal suspension (sol) and the

transition of the liquid “sol” into a wet and continuous network (gel). Removal of

the liquid from the sol yields the gel, and the sol-gel transition controls the

particle size and shape. Calcination of the gel then produces the oxide. The gel

phase can be processed by various drying methods to develop materials with

distinct properties. Subsequent drying under supercritical conditions converts the

gel into a low-density, highly porous aerogel, while drying induced by heating

typically results in a xerogel (low temperature) or a dense ceramic (high

temperature) (Figure 2.1)7.

41

Figure 2.1 Schematic representation of sol-gel process of synthesis of nanomaterials 7.

Two reactions are typically used in the sol-gel process: (1) there is an initial

hydrolysis reaction through which the alkoxide group (-OR) of the precursor is

replaced by the hydroxyl group (Equation 2.1). For example, the mechanism is

based on the nucleophilic attack to the central Si atom in silica production; (2)

this is then followed by water or alcohol condensation reactions (Equation 2.2

and 2.3), in which two hydrolyzed species (monomeric and polymeric silica

reacting units) link together to form siloxane bonds (Si-O-Si) with the elimination

of water or alcohol. Under most conditions, polycondensation commences before

hydrolysis is complete. However, conditions such as, pH, H2O/Si molar ratio, and

catalyst can force completion of hydrolysis before condensation begins 8.

Additionally, because water and alkoxides are immiscible, a mutual solvent such

as an alcohol is utilized9. With the presence of this homogenizing agent,

hydrolysis is facilitated due to the miscibility of the alkoxide and water.

42

Typical sol-gel processes require strong acid or base for accelerating the

hydrolysis of the precursors. It is generally found that the alkaline conditions

usually favor the formation of “particulate” sols, whereas the acidic conditions

produce weakly branched “polymeric sols”. For example, the kinetics and

mechanism of silica-particle formation by the base-catalyzed hydrolysis of TEOS

in alcohol media have been studied extensively10. It was found that the dilute

NaOH-catalyzed hydrolysis of TEOS had a first-order dependence on the

concentrations of both TEOS and hydroxyl ion (OH-). While for

ammonia-catalyzed reaction (Stöber process), both the rates of silica-particle

growth and TEOS hydrolysis were first order with almost the same specific rate

constant, indicating that silica-particle growth was reaction-controlled by the

hydrolysis of TEOS.

Typical precursors for the sol-gel synthesis of oxide materials include metal

alkoxides and metal salts11,12, among which the most versatile precursors are

undoubtedly alkoxides because they react readily with water. Alkoxide materials

consist usually of a metal or metalloid element surrounded by the reactive ligands.

Sol-gel methods using metal alkoxides usually produce fine and spherical oxide

particles of uniform size. However, the disadvantage of such water reactivity is

that tight control of the reaction conditions is required.

(2.2)

(2.1)

(2.3)

http://en.wikipedia.org/wiki/Precursor_(chemistry)

http://en.wikipedia.org/wiki/Metal

http://en.wikipedia.org/wiki/Alkoxide

43

The morphology and properties of a particular sol-gel inorganic network are

related to a number of factors that influence the rate of hydrolysis and

condensation reactions, such as, pH, temperature, reagent concentrations, alkyl

groups in the alkoxide, type of solvent, catalyst adopted and its concentration,

H2O/alkoxides molar ratio and drying 1, 9,13,14. Among the factors listed above, pH,

nature and concentration of catalyst, H2O/Si molar ratio, and temperature have

been identified as the most important.

2.2 CNT-Inorganic nanohybrids

2.2.1 Introduction to Carbon Nanotubes

2.2.1.1 Structures

Carbon nanotubes (CNTs) are the 1D allotrope of carbon and are formed by

predominantly sp2-bonded carbon atoms arranged in a honeycomb lattice. CNTs

are generally classified as either single-walled nanotubes (SWNTs) or

multi-walled nanotubes (MWNTs). A SWNT can be visualized as a single layer of

graphene sheet rolled into a seamless cylindrical tube with a diameter of 1–2 nm.

While a MWNT consists of several concentric and closed graphene tubules with

an overall diameter of ~10 to 100 nm and a length of up to centimeters. The

interlayer distance between the tubules is approximately 0.34 nm, similar to the

interlayer spacing in HOPG. Both types are displayed in Figure 2.2. CNTs can be

either open-ended or closed by a cap which in ideal models is described as a

hemispherical fullerene-type cap.

44

Figure 2.2 The structures of (a) SWNTs and (b) MWNTs 15.

2.2.1.2 Properties

Ever since their discovery in 1991 by Iijima 16, CNTs have drawn considerable

research attentions in the field of nanoscience and nanotechnology owing to their

rich electrical properties 17, high mechanical strength and excellent chemical and

thermal stability. CNTs possess high aspect ratio and large specific surface areas

attributed to their hollow geometry. While MWNTs are purely metallic, SWNTs

can be either metallic or semiconducting depending primarily on their diameter

and chirality.

The chirality is defined as the symmetry of a nanotube’s wall. A SWNT can be

considered as a rolled-up graphene sheet and is characterized by the way the

graphene sheet is conceptually rolled up to form it (Figure 2.3), i.e. the chiral

vector 18:

Ch=na1+ma2=(n, m) (2.4)

Where Ch is the chiral vector, a1 and a2 are unit vectors, n and m are integers

denote the number of unit vectors along two directions in the crystal lattice of

graphene. The length of Ch determines the tube diameter and the angle between

Ch and the (n,0) lattice vector, the chiral angle θ, determines the chirality.

(a) (b)

http://www.google.co.uk/url?sa=i&rct=j&q=&esrc=s&frm=1&source=images&cd=&cad=rja&docid=y0NOBr51zkX0dM&tbnid=UUSHfx1wLjTkfM:&ved=0CAUQjRw&url=http%3A%2F%2Fnanoparticles-environment.wikispaces.com%2F4.3.3%2BNano-filtration&ei=CB0PUcXAMabL0AWtyYHYDg&psig=AFQjCNHThzvXoNkkB_6swjncYyIz8bhLhg&ust=1360031321037024

http://en.wikipedia.org/wiki/Vector_(geometric)

http://en.wikipedia.org/wiki/Crystal_lattice

45

Tubes having n = m (θ= 30°) are called armchair NTs and those with m = 0 (θ= 0°)

are called zigzag NTs. Otherwise, they are called chiral NTs. Both armchair and

zigzag NTs have a high degree of symmetry. All the armchair tubes are metallic

and for zigzag and chiral NTs, when (n−m)/3 is an integer, the tubes are metallic

and otherwise semiconducting 19,20.

The situation in MWNT is complicated as their properties are determined by

contribution of all individual shells having different chiralities. However, it has

been reported for small diameter MWNTs that only one concentric tube needs to

be metallic for the overall electronic properties to be essentially metallic 21. In

large MWNTs quantum confinement is lost in the circumference.

Figure 2.3 Schematic representation of a 2D graphene sheet with the lattice vectors

a1 and a2 and the roll-up vector Ch=na1+ma2. The achiral cases, (n,0) zigzag and (n,

n) armchair are indicated with dashed lines. The translation vector T is along the

nanotube axis and defines the 1D unit cell. The shaded boxed area represents an

unrolled unit cell, defined by T and Ch. 18

46

2.2.1.3 Synthesis

There are three main methods for nanotube synthesis; electric arc discharge22,

laser ablation23, and chemical vapour deposition (CVD) 24. Although the former

two methods generally produce CNTs with fewer structural defects, they tend to

suffer from low yield issues and thus proves infeasible for mass production 25. On

the other hand, CVD shows great promise for possible industrial scaled-up due to

the relatively low growth temperature, high yields and high purities of the

synthesized CNTs 26. It is also capable of growing nanotubes directly onto the

desired substrate26, whereas the nanotubes produced from the other routes must

be subsequently processed and deposited in the required morphology. CVD

technique allows the growth of aligned CNTs of various packing densities which

may be useful for applications such as electrodes. Positional control of growth

has been achieved by patterned pre-deposition of the catalyst. In addition, this

method allows greater control over the morphology of CNTs by manipulating the

reaction parameters, such as reaction temperature, catalyst concentration and

reaction time 26-28.

In CVD technique, the nanotubes are grown from carbon containing gaseous

compounds (i.e. hydrocarbon) which are reacted with a metal catalyst at moderate

temperatures (≤ 1000 °C). The catalyst is present either in-situ from a precursor

or pre-produced on a substrate. However, this method is not without drawbacks.

Residual metal catalyst particles tend to remain in the CNT structures which limit

some of their applications, therefore post-production treatments are required to

purify the nanotubes.

2.2.1.4 Applications

The unique physical and chemical properties of CNTs have led to their diverse

use as supercapacitors29, reinforcement materials of polymers and ceramics 30-32,

electromechanical actuators 33, field emission devices 34, gas sensors 35,36 and

nanosize probe tips for AFM37. Recently their bioapplication as biosensor

47

materials has attracted increasing interests due to their ability to enhance the

electroactivity of biomolecules and to promote the electron-transfer between the

biomolecules’ active site and the electrochemical transducer.38-41 CNTs can also

act as supports for metal and semiconductor catalysts thanks to their high aspect

ratio 42-45.

To take advantages of the remarkable properties of CNTs, a popular solution is to

prepare composite materials based on CNTs and various other materials ranging

from ceramics, polymers to biomolecules 46. However, the as-produced CNTs

tend to be chemically inert due to their inherently hydrophobic nature which

provides little attractive interaction with the inorganic compounds. Thus, it is

necessary to modify their surface chemistry in order to achieve good interfacial

bonding with the matrix in the composites.

2.2.2 Functionalization of CNTs

Due to the hydrophobic nature of pristine CNTs, they tend to aggregate into

bundles in solvents held together by the strong van der Waals forces. This

bundling is a significant barrier to their processing and also perturbs the

electronic structure of the tubes. Functionalizaiton of CNTs has opened up the

possibilities of dispersing 47,48 and self-assembly of the nanocarbons49 which

allows for the generation of useful architectures50.

Two strategies are generally reported towards the functionalization of CNTs: (1)

covalent functionalization through attachment of chemical groups to the sidewall

of CNTs51,52 and (2) non-covalent adsorption of various functional molecules,

such as surfactants and polymers. Both non-covalent and covalent

functionalization have been reported to improve the solubility of CNTs 53 which

is necessary for their characterization and manipulation.

Recently, increasing interests have been focused on the functionalization of CNTs

48

with biomolecules as motivated by the prospects of using nanotubes as new types

of biosensor materials 54-56. Carbohydrates57, proteins 55,56, peptides58,59, and

single strand DNA55 have been demonstrated to modify CNTs through either

covalent 60 or non-covalent way49,61-63. The modification of nanotubes by these

biomolecules, as well as their analogs and precursors (such as oligosaccharide,

amino acids and peptide, etc.) represents a significant step toward the application

of CNTs in the field of biotechnology and the transfer of biomolecular

self-assembly techniques to nanomaterials.

2.2.2.1 Covalent functionalization

Covalent functionalization of CNTs provides more control over the location and

density of the attached groups than the non-covalent adsorption and thus leads to

more robust and predictable conjugates. Considerable progress has been made on

the open-end 64,65 and sidewall 52,66 modifications of CNTs using covalent

chemistry. Reactions that are usually employed to introduce chemical

functionalities onto CNTs include cycloadditions 67,68 nucleophilic additions 69,

ozonolysis70, halogenation 71 and radical additions 72-74.

The most popular example of covalent functionalization involves the oxidation of

CNTs in strong acids, such as HNO3 75,76 and HNO3/H2SO4 mixture 64,77. Acid

treatment introduces oxygen-containing functional groups (-COOH and –OH)

onto the sidewalls and open ends of CNTs 76,78 which significantly enhance their

aqueous solubility and facilitates further functionalization 79-83. Fu et al. 84 have

developed a milder route for attaching bovine serum albumin (BSA) protein to

CNTs via the esterification of nanotube-bound carboxylic acids by oligomeric

polyethylene glycol compounds followed by the ester-to-amide transformation

reactions with BSA protein. The entire conjugation procedure did not subject the

protein to any damaging experimental conditions, therefore, the method may be

valuable for the preparation of conjugates involving more fragile biological

species.

49

However, this method suffers from a major disadvantage of cutting the CNTs in

short lengths, making them useless for some applications. Soluble full-length

CNTs have been reported using the 1,3-dipolar cycloaddition of an

aminoethylene glycol linker to the external surface of CNTs and the

derivatization with N-protected glycine was then obtained via amidation reaction

(Figure 2.4). 85

Figure 2.4 1,3-dipolar cycloaddition of an aminoethylene glycol linker to the

external surface of CNTs and the derivatization with N-protected glycine was then

obtained via amidation reaction. 85

Biofunctionlization of CNTs has also been employed in the fabrication of

nanoscale biosensors base on enzyme-CNT86, DNA-CNT87 or antibody-CNT

conjugates88. Yu et al. 41 attached myoglobin and horseradish peroxidase

covalently onto the ends of vertically oriented SWNTs forest arrays, which were

used as electrodes. Results suggested that the “trees” in the nanotube forest

behaved electrically similar to a metal, conducting electrons from the external

circuit to the redox sites of the enzymes.

Lin et al. 89 have demonstrated a novel glucose biosensor based on CNT

nanoelectrode ensembles (NEEs) for the selective and sensitive detection of

glucose. Glucose oxidase (GOx) was covalently immobilized on CNT NEEs via

carbodiimide chemistry by forming amide linkages between their amine residues

and carboxylic acid groups on the CNT tips (Figure 2.5). The biosensor

50

effectively performed a selective electrochemical analysis of glucose in the

presence of common interferents (e.g., acetaminophen, uric and ascorbic acids),

avoiding the generation of an overlapping signal from such interferers. Such an

operation eliminates the need for permselective membrane barriers or artificial

electron mediators, thus greatly simplifying the sensor design and fabrication.

Figure 2.5 Fabrication of a glucose biosensor based on CNT nanoelectrode

ensembles: (A) Electrochemical treatment of the CNT NEEs for functionalization

(B) Coupling of the enzyme (GOx) to the functionalized CNT NEEs. 89

Covalent functionalization has been shown to introduce structural defects to

CNTs’ sidewall which lead to a disruption of the nanotubes’ delocalized π system

and consequently compromises their electronic and mechanical properties 90. This

will in turn lead to a significantly poorer performance of CNT-based composites.

To circumvent this problem, non-covalent modification of CNTs which do not

significantly alter their properties is required for the development of high

performance CNT-based hybrids. Nevertheless, the covalent route offers a

51

convenient and controllable means of tethering molecular species.

2.2.2.2 Non-covalent functionalization

In contrast to covalent functionalization, non-covalent binding, which utilizes π-π

stacking,62,91,92 hydrophobic interactions93,94, electrostatic interaction 95 and

hydrogen bonding has relatively less impact on the structural and functional

properties of CNTs.

Stable CNT dispersions in both aqueous and organic solutions have been

achieved through immobilization of ionic 96-98 and nonionic surfactants 99,100

respectively. This solubilization process opens the door to solution chemistry on

pristine CNTs. Commonly employed surfactants include the anionic sodium

dodecyl sulfate (SDS) 101-104 and sodium dodecyl benzene sulfonate (SDBS) 105,106,

cationic CTAB98 and nonionic surfactants such as Triton X-100 and Tween 80.

Synthetic aromatic ligands such as pyrenyl 107-112, porphyrin 113 as well as phenyl

groups are known to interact strongly with the sidewalls of CNTs via π-π stacking

interaction. The interaction is typically weaker for phenyl groups as compared

with the ligands with higher aromaticities. However, their smaller size favors the

higher density of the groups that can be immobilized on CNTs. Dai et al. 62 have

reported the non-covalent functionalization of the sidewalls of SWNTs with a

bifunctional molecule, 1-pyrenebutanoic acid, succinimidyl ester (Figure 2.6,

molecule 1), and subsequent immobilization of various biological molecules onto

nanotubes with a high degree of control and specificity.

52

Figure 2.6 Amine groups on a protein react with the anchored succinimidyl ester to

form amide bonds for protein immobilization.62

There has been increasing interest on using biologically based surfactants, which

then opens up the biochemistry tool kit to nanotechnology. Biomolecules such as

DNA112, polysaccharides57,114 and peptides62,91,92 have been reported for

functionalization of CNTs. Among the biosurfactants, peptides are of particular

interest owing to their designable chemistry. Phage display study has identified

the CNT binding peptide sequences which were invariably rich in aromatic amino

acids such as histidine (H) and tryptophan (W), with W, in particular, interacting

strongly with the nanotube surface 115. It was suggested that the aromatic rings in

these amino acid residues contributed to the observed affinities through π-π

stacking interactions92. However, these literatures tend to focus on the long chain

peptides which usually contain 12 or more residues that are expensive to produce.

As a cost-effective alternative, recent studies have reported using synthetic

aromatic ligands combined with aromatic amino acids for CNT dispersion.

Cousins et al. 47 have demonstrated the use of N-(fluorenyl-9-methoxycarbonyl)

terminated aromatic amino acids (Fmoc-AAs) as surfactants for preparing

homogeneous CNT dispersions (Figure 2.7a). (It should be noted that the author

of this thesis was a co-author on this paper.) Fmoc was selected as a particularly

promising ligand since it is used commonly as a protecting group in solid-state

53

peptide synthesis and it is known to be able to self-assemble into nanofibres via

π-π stacking interactions 116. The turbidity study of the dispersions of CNTs in the

Fmoc-AA solutions revealed the comparable ability of these biosurfactants to

disperse CNTs to those achieved by using commonly used surfactants such as

SDS and SDBS. The molecular interactions between the ligand and nanotube

surface were then confirmed by quantum mechanical modelling and it was found

that both the aromatic fluorenyl rings and the aromatic rings in the side chains of

the amino acid were stacked on the surface of CNTs to maximize their π-stacking

interactions (Figure 2.7b).

(a) (b)

Figure 2.7(a) TEM micrographs of MWNTs dispersed with Fmoc-W (trp). Arrows

indicate the edge of the lattice structure upon which Fmoc-W aggregates are

apparent; (b) Optimized structures of (i) Fmoc-G (gly) and (ii) Fmoc-W bound to

[6,6] SWNTs with close-up images that highlight the orientation and arrangement

of Fmoc and the aromatic W ring 47.

Although the approaches described above increase the solubility of CNTs, they

have not been generally adapted to control the assembly of the solubilized CNTs

into higher order architectures that are necessary for realizing many of their

applications. Dieckmann et al.49,50 have designed an amphiphilic α-helical peptide

(“nano-1”) not only to coat and solubilize CNTs into water, but also to control the

self-assembly of the peptide-coated nanotubes into supramolecular structures

through peptide-peptide interactions between adjacent peptide-wrapped

nanotubes. The CD measurements suggested that the α-helical conformation of

the peptide is stabilized in the presence of the nanotubes through the interaction

of the hydrophobic face of the helix with the nanotube surface. Electron

54

microscopy and polarized Raman studies revealed that the peptide-coated

nanotubes assemble into fibres with the nanotubes aligned along the fibre axis.

Most importantly, the size and morphology of the fibres can be controlled by the

addition of either salt in different concentrations or the amphiphilic additive DMF

which can affect the peptide-peptide charge interactions (Figure 2.8). This study

helps to realize the transfer of biomolecular self-assembly techniques to

nanomaterials.

(i) (ii)

55

Figure 2.8 (i) SEM images of nano-1/SWNT fibres formed from a 100 μM

peptide/nanotube dispersion upon addition of no salt (A), 40 mM NaCl (B), and 120

mM NaCl (C). (ii) (A) SEM image of fibres formed from the addition of 0.0015%

(by volume) DMF to a nano-1/SWNT dispersion. (B) Low-resolution TEM image of

the same fibres observed in i(A). The small dark spheres are Fe catalyst particles

from the HiPco SWNT synthesis. (C) High-resolution TEM image of the same fibres

showing alignment of nanotubes. The large dark areas are Fe particles 49.

These investigations have contributed to the understanding of the nonspecific

interactions between CNTs and biomolecules, and the current knowledge on

non-specific protein–nanotube interactions has already been applied to the

development of biosensors but they have also revealed the complexity of the

issue. Researches based on the molecular level are required to further understand

the interactions.

Polymer wrapping has also been reported for CNT dispersion without destroying

their electrical character117,118. The wrapping of SWNTs with polymers that bear

polar side-chains, such as polyvinylpyrrolidone (PVP) or polystyrenesulfonate

(PSS), leads to stable solutions of the corresponding SWNT/polymer complexes

in water 117. The thermodynamic driving force for complex formation is the need

to avoid unfavorable interactions between the apolar tube walls and water. It is

thought that multi-helical wrapping of the tubes with the polymers is most

favorable for reasons of strain. A nonionic surfactant or polymer’s ability to

suspend nanotubes appears to be due mostly to the size of the hydrophilic group,

with higher molecular weights suspending more nanotube material because of

enhanced steric stabilization with longer polymeric groups119.

An “unzipping” mechanism for nanotube isolation from a bundle with the

combined assistance of ultrasonication and surfactant adsorption has been

proposed as shown in Figure 2.9120. The role of ultrasonic treatment is likely to

provide high local shear, particularly to the nanotube bundle end (ii). Once spaces

or gaps between the bundle and individual nanotubes at the bundle ends are

formed, they are propagated by surfactant adsorption (iii), ultimately separating

the individual nanotubes from the bundle by either steric stabilization or

56

electrostatic repulsions (iv).

Figure 2.9 Proposed mechanism of nanotube isolation from bundle (i) obtained by

ultrasonication and surfactant stabilization. Ultrasonic processing “fray” the

bundle end (ii), which then becomes a site for additional surfactant adsorption. This

latter process continues in an “unzippering” fashion (iii) that terminates with the

release of an isolated, surfactant-coated NT in solution (iv).120

Several mechanisms have been proposed for the stabilization of CNT dispersion

by surfactants. O'Connell et al. 96 have suggested the formation of SDS

cylindrical micelles around SWNT (Figure 2.10a) or the hemimicellar adsorption

of the surfactants on the tubes (Figure 2.10b) while Richard et al. 121 suggested

the formation of helices or double helices, and Yurekli et al.101 suggested that the

structureless random adsorption with no preferential arrangement of the head and

tail groups of the surfactants is responsible for the stabilization of the dispersions

(Figure 2.10c).

57

Figure 2.10 Schematic representations of the mechanisms by which surfactants help

disperse SWNTs. (a) SWNT encapsulated in a cylindrical surfactant micelle: right:

cross section; left: side view. (b) Hemimicellar adsorption of surfactant molecules

on a SWNT. (c) Random adsorption of surfactant molecules on a SWNT.101

2.2.3 CNT-inorganic nanohybrids

During the past decades, CNT based hybrid materials have been extensively

reported owing to their potential in applications such as photocatalysis122,123,

electrocatalysis124-127, gas and biosensing128-131, supercapacitors132-135 and field

emission device 136-141.

The first CNT based nanohybrid was produced by opening the capped tube ends

of MWNTs and then filling the hollow cavities with lead particles 142. Later,

SWNTs were filled with RuCl3 143. Although a wide range of compounds have

been successively encapsulated into both SWNTs and MWNTs, few have

exploited their potentials in application and have been mainly used by electron

microscopists to understand crystallization in restricted volumes.

58

Alternatively, a wide range of inorganic compounds have been anchored onto the

surface of CNTs for the preparation of hybrid materials. Among the inorganic

components, the most frequently studied are semiconductor oxide nanoparticles

such as SiO2144-147, Al2O3

148-150, SnO2151-153, ZnO154-156 and TiO2

157-163. Of

particular interest are dielectric materials such as silica and TiO2. TiO2 exists in

nature as three polymorphic forms, namely rutile, anatase and brookite, amongst

which, the most important being rutile and the metastable anatase phases. Both of

the phases have tetragonal structures. The properties and applications of TiO2 are

greatly dependent on their crystalline phase, particle size, and morphology, which

could be controlled by varying the reaction conditions 164. A number of studies

have reported the improved photocatalytic activity of CNT-TiO2 hybrids as

compared to the individual component for the oxidative degradation of organic

compounds 122,123,165.

2.2.3.1 Synthesis

The most important challenge in synthesizing such hybrid materials is optimizing

the interface between CNTs and the inorganic components. In general, two

strategies have been adopted for the synthesis of CNT-inorganic hybrids;

1. ex-situ techniques where the preformed inorganic components are

directly attached to the surface of CNTs,

2. in-situ techniques where the inorganic components form directly on the

surface of pristine or functionalized CNTs.

The ex-situ route is mainly used for the deposition of metal nanoparticles67,166 and

semiconductor QDs 167. Surfactants 168 are usually employed as the linking agents

in this approach which utilize both covalent 67,167,169 and non-covalent

interactions 169-176.

Although the ex-situ route holds the advantage of producing inorganic

components with desired structures and dimensions, it requires the chemical

modification of either CNTs or inorganic compounds for their attachment.

http://en.wikipedia.org/wiki/Rutile

http://en.wikipedia.org/wiki/Anatase

http://en.wikipedia.org/wiki/Brookite

59

Furthermore, the in-situ route allows more flexibility of the morphology of the

deposited inorganic components as either discrete units in the form of

nanoparticles or a continuous film on CNTs, while the ex-situ way is typically

restricted to the formation of monolayers of nanoparticles. The presence of CNTs

also prevents the growth of crystals during crystallization and

phase-transformation thus provides an efficient way of synthesizing nanohybrids

with high specific surface area.

The in-situ techniques include (1) hydrothermal techniques132,148,155,159 (2) sol-gel

process 177 (3) electrochemical methods 178-181 and (4) gas phase deposition 182-184.

The main advantage of hydrothermal technique is that it enables the formation of

crystalline phase without the need for post-annealing and calcinations. However,

it typically requires high temperatures 132,148,155. Jitianu et al. 159 have compared

the morphologies of TiO2 coating on CNTs obtained from both sol-gel process

and hydrothermal methods and found that the coating produced with

hydrothermal method is less uniform and the nanotubes surface is partially

damaged due to the oxidizing medium of deposition.

To overcome the above problems, sol-gel process has been widely employed as

an alternative method to prepare CNT-inorganic nanohybrids benefiting from its

benign reaction conditions. Sol-gel process on both covalently152,161,185-188 and

non-covalently147, 189, 190 functionalized CNTs have been reported. During the

sol-gel process, CNT surface chemistry plays an important role in inducing

inorganic compound deposition as well as in controlling the structure and

property of the deposited coatings172.

2.2.3.2 CNT-SiO2 hybrids

Silica-CNT hybrids are of great interest due to their potential in the development

of nanoscale sensors and electric devices 191,192 as well as optical, magnetic, and

catalytic applications 193. CNT-SiO2 hybrids also combine the bioactivity of silica

and the conductivity of CNTs which facilitate their biomedical applications 194. In

60

particular, they can promote compatibility with existing silicon based technology.

CNTs have been widely used as reinforcing fillers for silica. However, due to

poor interfacial bonding of pristine CNTs with silica matrix, they are very easily

agglomerated in silica matrix, which greatly inhibits the effective load transfer.

Previous report has demonstrated the covalent functionalization of CNTs with

silane followed by covalent incorporation of the functionalized CNTs into the

silica matrix via the sol–gel process 32. More uniform dispersion of CNTs as well

as stronger interfacial interaction between CNTs and the matrix were achieved,

thus lead to improved mechanical properties and higher electron-transfer kinetics.

But for electronic applications, the covalent functionalization would inevitably

lead to disruption of the nanotubes’ delocalized π system. Alternatively, by

employing CNT-SiO2 nanohybrids as reinforcing filler, the unique electronic

property of CNTs was greatly preserved 32,195.

Seeger et al.144 have employed PEI as cationic surfactant to coat MWNTs with

SiOx. The adsorption of PEI creates positive charges on CNTs’ surface and thus

favors the deposition of negatively charged colloidal SiOx. The same author also

studied the interface between CNT and SiOx coating in a bulk composites

produced using sol-gel method. HRTEM analysis suggested that the outer

nanotube shells are strongly bonded to the SiOx matrix due to the carbothermal

reduction. Theoretical models accounting for the stable SiOx/CNT interfaces were

further proposed using density functional based tight binding method (DFTB) 145.

Bourlinos et al. 147 have prepared water-dispersible CNT–silica hybrids by

wetting the hydrophobic surface of pristine CNTs with vinylsilane molecules

through the non-covalent interactions between the vinyl groups and the aromatic

walls of CNTs. Subsequent condensation of silane leads to bonded oligomeric

siloxane species that upon calcination afford ultrafine silica nanoparticles

embedded in CNTs.

61

Whitsitt et al. 146 have investigated the effect of surfactants on the morphology of

silica coated SWNTs by employing two kinds of surfactants, i.e. the anionic

surfactant SDS and the cationic surfactant DTAB. It was found that the use of

SDS result in the formation of coated ropes while DTAB lead to individually

coated SWNTs. The author attributed this effect to the pH stability of the

surfactant-SWNT interaction. This paper also indicates that the coating does not

alter the electrical properties of SWNTs. Furthermore, the SiO2 coating allows

ready etching to selectively expose CNTs for applications in addressable SWNT

devices and sensor.

Li et al. 196 have reported a non-covalent strategy for the production of CNT-SiO2

nanohybrids based on the π-π stacking of a bifunctional molecule 1-aminopyrene

on CNTs (Figure 2.11a). Subsequently, the amino groups present on CNTs’

surface specifically adsorb silica precursors via electrostatic interactions (b)

followed by the in-situ formation of silica nanoparticles (c).

62

Figure 2.11. Scheme for the preparation of CNT–silica nanohybrids.196

Silica NTs were successfully fabricated from calcination of SiO2-MWNT hybrids

which was obtained through a covalent approach 197. As shown in Figure 2.12,

MWNTs were first oxidized in KMnO4 followed by grafting of a silane coupling

agent AEAPS. The amine groups in grafted AEAPS on MWNTs could activate

the silica shell formation by acid–base interaction. The oxidation of CNTs was

performed in the presence of a phase transfer catalyst which helps to minimize

the damage to CNTs during the oxidation process 198 as well as to improve the

selectivity of the functional groups introduced on CNTs.

63

Figure 2.12 Scheme of the reaction between MWCNT-OH and AEAPS for the

following synthesis of silica coated MWCNTs. 197

Fu et al. 199 have used APTES (3-aminopropyltriethoxyysilane) as a promoter to

coat pristine SWNTs with a thin layer of SiO2. APTES could be adsorbed onto

CNTs through the interaction between the amine groups and SWNTs sidewalls.

Upon heating the adsorbed APTES polymerized to form a primer layer onto

which a thin layer of SiO2 is grown through a modified Stober method. The rich

chemistry of silica allows for further functionalization of SWNTs with a large

variety of functional groups for applications as highly sensitive gas sensors.

Satishkumar et al. 200 reported the direct coating of SiO2 on both pristine and

acid-treated CNTs. It was found that the acid-treated CNTs generally give better

oxide coating over the pristine CNTs. SiO2 NTs containing transition metal ions

have also been prepared which show potentials in catalysis.

2.2.3.3 CNT-TiO2 hybrids

TiO2 is of particular interest due to its excellent photocatalytic activity, relative

non-toxicity, and long-term thermodynamic stability. It has shown great potential

in the field of environmental protection, including water and air purification 201-203,

gas sensing 204-206 and dye-sensitized solar cells (DSSCs) 207,208. Following on

from the pioneer work on the photocatalytic water splitting reported by Fujishima

64

and Honda in 1972 209, the photocatalytic properties of TiO2 have been widely

used to convert solar energy into chemical energy in the form of hydrogen 210, 211

and hydrocarbons 212. The photocatalytic degradation of organic pollutants such

as phenol 213, acetaldehyde 214 and methylene blue dye 215 by TiO2 has also been

extensively studied.

The photocatalytic properties of TiO2 are derived from the formation of

photogenerated electrons and holes which occurs upon irradiation of UV light

with energy greater than the band gap of TiO2 (Figure 2.13)216 .The

photogenerated holes in the valence band diffuse to the TiO2 surface and react

with the adsorbed water molecules, forming hydroxyl radicals (•OH). The

photogenerated holes and the hydroxyl radicals oxidize nearby organic molecules

on the TiO2 surface. Meanwhile, the photogenerated electrons in the conduction

band participate in the reduction processes, which are typically reacted with

molecular oxygen in the air to produce superoxide radical anions (O2− •).

Figure 2.13 Mechanism of photocatalysis on the surface of TiO2 in presence of UV

radiation. 216

Factors incuding specific surface area, pore volume, pore structure, and

crystalline phase could significantly affect the photocatalytic performance of

TiO2. Nanostructured TiO2 materials offer higher specific surface area and show

enhanced sensitivity, electronic conductivity and photovoltaic activity. Both

anatase and rutile phases have been widely used as white pigments with anatase

showing greater potential in photocatalysis due to its higher photocatalytic

65

activity. Whilst rutile phase possesses a higher absorbance property and refrective

index which facilitate their use as sunscreens.

One of the widest uses of TiO2 in functional applications is in dye sensitized solar

cells (DSSCs). DSSCs are based on the photo-injection of electrons from dye

molecules into a wide band gap semiconductor and holes transport by a redox

electrolyte. Until now the most efficient DSSCs are still based on TiO2 electrodes.

The anodes of a typical Grätzel-type DSSC are usually constructed from

nanoporous dye-sensitized nanoparticulate TiO2 films (Figure 2.14) 217 which

could drastically enhance the solar energy conversion efficiency due to their high

surface area 218. Currently, the highest power conversion efficiency achieved for

DSSCs is up to 12% 219.

Figure 2.14 Schematic representation of a dye-sensitized solar cell based on

particulate TiO2. 217

However, the nanoparticulate structure is not optimized to carry electrons from

the injection site at the particle surfaces towards the anode. In contrast, the 1D

66

nanostructures, especially the nanotubular structures, allow for direct conduction

paths (Figure 2.15) 220,221 which is favorable for photocatalytic reactions.

Figure 2.15 Schematic representation of the electron path through a (a) percolated

and (b) oriented nanostructure. 220

CNTs act as good support materials for photocatalyst particles due to their high

mechanical and chemical stability as well as their mesoporous nature which

favors the diffusion of reacting species.

Taking advantage of the superb conductivity of CNTs, we may expect that the

combination of CNTs with TiO2 may produce interesting charge-transfer

behavior and further enhance the photocatalytic activity of TiO2. A considerable

progress has been made so far in the fabrication and application of these

promising hybrid materials and a few examples are worth noticing. For instance,

by coupling SWNTs with TiO2 in a photovoltaic device, the photo-conversion

efficiency was largely enhanced from 7% for the pure TiO2 to 15% for the hybrids 222.

This was attributed to the beneficial role of SWNTs as conducting scaffold to

collect and direct the flow of photogenerated electrons in nanostructured

semiconductor films thus considerably lower the probability of recombination

(Figure 2.16).

(a) (b)

67

Furthermore, the biocompatibility of CNTs could be improved by constructing

CNT-TiO2 nanohybrids, and the electrodes modified with the nanohybrids show

great potential in the development of biocompatible and multi-signal responsive

biosensors for the early diagnosis of cancers223.

Huang et al. 162 have reported the immobilization of rutile TiO2 on the sidewall of

acid-treated MWNTs which exhibited three distinct morphologies depending on

the reaction temperatures. They further demonstrated that the crystal structure of

CNTs may suffer some damage during the harsh acid treatment. Furthermore,

such treatment provide relative poor control over the density, type, and location

of the introduced functional groups which in turn lead to the non-uniform coating

of inorganic compounds. Therefore, acid treatment of CNTs should be avoided in

order to preserve their superior electronic properties. Nevertheless, this method

did lead to the mechanically stable coating. Alternatively, many studies have

utilized non-covalent adsorption of surfactants through interactions such as π-π

stacking196 and hydrophobic interactions190 to alter the surface chemistry of CNTs

for preparing CNT-inorganic hybrids.

Eder et al. 172 have employed benzyl alcohol (BA) as a surfactant to coat pristine

CNTs with TiO2 without the need of covalent functionalization. They assume that

the benzene ring of the surfactant adsorbed on CNT surface via - stacking

interactions, while simultaneously providing hydroxyl groups for the hydrolysis

Figure 2.16 Electron transport across nanostructured semiconductor films: (A) in

the absence and (B) in the presence of CNTs support. 222

68

of the titanium precursor (Figure 2.17Left). The hypothesis is further confirmed

by the molecular dynamics simulation 224 which showed that the phenyl rings in

the BA favour an orientation parallel to the graphene sheet enabling π-π stacking.

Furthermore, the presence of BA was also found to strongly affect both the

crystallization and phase transformation by providing very small and uniform

anatase and rutile nanocrystals (Figure 2.17Right). However, BA also tended to

retard the phase transformation due to the formation of a carbon coating on TiO2

upon heating in argon. They further showed that the TiO2 coating catalyzed the

oxidation of CNTs via a Mars–van Krevelen mechanism, thus reducing the

required oxidation temperature by about 120 °C 160.

Figure 2.17. Left: Scheme of the beneficial role of benzyl alcohol (BA) in the in situ

coating of pristine CNTs with TiO2. One possible conformation of two BA molecules

on the CNT surface is shown in Scheme. Right: SEM images of TiO2 on CNTs after

conversion from anatase to rutile: A) no BA and B) with BA. 172

Both silica and titania coatings on SWNTs were achieved by templating a

multifunctional peptide P1R5 modified SWNTs, where the R5 peptide repeat unit

of silaffin that is involved in diatom biosilicification is capable of precipitating

silica from the hydrolyzed silica precursors at room temperature 91. The dual role

of the peptides has been demonstrated as: (1) to not only suspend SWNTs in

solution (2) but also direct the deposition of silica and titania at the surface of

CNTs. Electron microscopy analysis showed that SWNTs were embedded within

silica and TiO2 matrix rather than individually coated, and that the precipitation

reaction led to agglomeration of SWNTs.

69

2.2.3.4 Inorganic nanotubes

The inorganic NTs can be prepared by calcination of the corresponding

CNT-inorganic nanohybrids to remove CNT templates 200. The major advantage

of template-assisted sol-gel technique is that it allows ready control over the

dimension of the synthesized inorganic NTs by the size of the CNT templates and

by adjusting the reaction conditions, such as reaction time 151, reaction

composition, and the choice of metal precursor 160.

Previous study has demonstrated that the presence of inorganic coating could

affect the oxidation stability of CNTs, which can be utilized for the production of

inorganic nanotubes. A wide range of oxide NTs have been produced using CNTs

as templates, including V2O5 200,225, SiO2

197,200, Al2O3 200, RuO2

226, Fe2O3 227 and

TiO2 160, 228, 229. Their tubular morphology has facilitated their application in gas

sensing and photocatalysis.

Among the inorganic NTs, TiO2 nanotubes are of special interest. Several recent

studies have demonstrated the improved properties of TiO2 NTs compared with

colloidal and other forms of TiO2 for applications in photocatalysis 155, 230-232,

sensing233, dye-sensitized solar cells 234 and photovoltaics 234. Their improved

performance was generally attributed to the high specific surface area 231,235, the

high electron transfer 234 and ready access to the reactants provided by the tubular

structures 229. However, Bouazza et al. 236 have demonstrated the higher

photocatalytic activity of CNT-TiO2 hybrids over TiO2 NTs for the oxidation of

propene at low concentration which was attributed to the synergistic role of

CNTs.

Adachi et al. 237 have reported the synthesis of single-crystalline TiO2 nanotubes

using templates formed by LAHC/TIPT with an ACA system. The dye-sensitized

solar cells based on these NTs showed more than double the short-circuit current

density than those made of TiO2 nanoparticles in the thin-film thickness region.

This was attributed to the high electron transfer through single-crystalline TiO2

70

nanotubes compared to that through nanoporous TiO2 films composed of

nanoparticles. The photoconversion efficiency of the TiO2 nanotube cells was

around 5%. They also showed the highest photocatalytic activity compared to the

commercially available nanocrystalline TiO2 particles 234.

TiO2 NTs can also act as supports for metal nanoparticles 229, which shows great

potential for both sensing 238 and catalytic applications 239. Grimes et al. 233

reported the use of TiO2 NTs coated with palladium particles lead to a 1000 fold

higher sensitivity for sensing small H2 concentration in the atmosphere, compared

to particle-based substrate. Not simply surface area but the NT morphology that is

responsible for this unprecedented sensitivity. Meng et al. 240 have employed

Pt-TiO2 nanotube hybrid electrodes for methanol oxidation which showed a much

enhanced electrocatalytic activity compared with that of pure Pt electrode. Su et

al. 241 studied the effect of the structure on the photocatalytic activity of Pt-doped

TiO2 nanotubes. They showed that Pt-doped TiO2 nanotubes with longer length

exhibited higher photocatalytic activity for the degradation of methyl orange

under UV and visible irradiation. Yu et al. 242 also demonstrated the potential of

TiO2 nanotube-supported Pt as catalyst for CO2 recycling and methane

production. Sato et al. 243 investigated the catalytic behavior of Pt/TiO2-NTs for

the water gas shift as well as the hydrogenation of CO. The observed catalytic

activities were two to three times higher compared to the conventional

impregnated Pt/TiO2 catalysts.

Vertically aligned TiO2 NT arrays have been demonstrated to significantly

accelerate the growth rate of osteoblast cells 244. Enzymes have been successfully

immobilized to the inside walls of the TiO2 NT arrays and it was shown to retain

its catalytic activity for a minimum of 4 days 245. The highly aligned TiO2 NT

arrays were initially synthesized through hydrothermal treatment of TiO2 particles

in concentrated NaOH solution at temperatures above 110 °C 246-248. However, the

strongly basic condition as well as the elevated reaction temperature are not

environmental friendly. Therefore, methods such as anodization of titanium films

71

249-252 and sol–gel templating of porous alumina membrane253-256 were developed

for the fabrication of highly ordered TiO2 NT arrays.

The anodization route requires specialized setup and cost extra energy. The length

of the resultant nanotube array was limited to ~500 nm in the aqueous electrolyte

249 and significantly longer tubes were achieved in the non-aqueous electrolytes

257-259. However, the synthesized NT exhibited an uneven surface with rings

present on the outer walls periodically 233, 249 and extra chemical etching is

required to remove the barrier layer 260.

On the other hand, the template-directed routes have the advantage of producing

TiO2 NT arrays with tunable diameter and length compared with the

electrochemical anodization. Sander et al. have prepared aligned TiO2 NT arrays

with well controlled dimensions by atomic layer deposition (ALD) of TiO2 within

a porous alumina membrane. The production of porous alumina membrane

templates requires anodization of Al film followed by chemical etching 261.

Besides, the TiO2 overlayer must be removed by mechanical polishing prior to

etching away the alumina template 253. Another limitation reported for such

template is that due to the weak driving force, direct filling of the nanosized pores

is very difficult 262. Therefore, such method failed to produce nanofibers with

diameters <50 nm. To overcome this problem, Miao et al. 263 employed an

electrochemically induced sol-gel method to prepare TiO2 single-crystalline

nanowire arrays with diameters ranging from 10 to 40 nm.

Sol-gel template method also allows the growth of superlong NTs which is key to

achieve improved photocatalytic efficiency 264. Zhao et al.265 reported the

production of millimeter long aligned TiO2/CNT arrays via electrodeposition

using superlong CNTs as templates which were prepared by CVD method. The

resultant TiO2/CNT coaxial arrays exhibited minimized recombination of

photoinduced electron–hole pairs and fast photocurrent responses. This study

establishes the base for the synthesis of super-long TiO2 NT arrays.

72

2.3 peptide self-assembly and mineralization

2.3.1 Introduction

Molecular self-assembly is a powerful “bottom-up” approach for fabricating

nanoscale architectures, for example in nature, ribosomal proteins and RNA

coalesce into functional ribosomes. Molecular self-assembly is mediated by weak,

non-covalent bonds-notably hydrogen bonds, ionic bonds (electrostatic

interactions), hydrophobic interactions, van der Waals interactions and

water-mediated hydrogen bonds. Amongst all the self-assembling biomolecules,

peptides represent the most favorable building blocks for the design and synthesis

of nanostructures because they offer a great diversity of chemical and physical

properties.

2.3.2 Strategies for peptide self-assembly

There are two general categories into which designed self assembling peptide

systems fall, natural and non-natural. The first category utilises the basic

conformational units of naturally existing proteins, β-sheets and turns, α-helices

and coiled coils. Through the examination of protein sequences it has been

possible to derive simple rules that promote the formation of one of the basic

conformational units. The second category covalently links amino acids to other

molecules; either an alkyl chain to form a peptide amphiphile (PA), or to an

aromatic group to create π–π interactions between the aromatic groups.

2.3.2.1 β-sheets and α helices

β-sheets consist of multiple peptide chains that have an extended backbone

connected laterally by hydrogen bonds between the backbone amides and

carbonyls. β-sheets can be orientated so that all of the N-termini of successive

strands are oriented in the same direction, described as a parallel structure, or so

that the successive β strands alternate directions, described as an anti-parallel

73

structure. This has an important impact on the orientation of hydrogen bonds

between sheets and side chain orientations and interactions. For example, Aggeli

et al. 266 reported the hierarchical self-assembly of chiral rod-like molecules as a

model for peptide β-sheet tapes, ribbons, fibrils, and fibres, all of which vary in

the number of sheets that pack together to form the final structure.

Different from “infinitely” assembling systems such as self-assembled

monolayers (2D) and nanofibres (1D), finitely assembling systems in which all

dimensions of assembly are controlled can be prepared through the utilization of

“molecular frustration” in which two or more components of the assembler have

opposing preferences for solvation environment, attraction versus repulsion, or

compact and ordered packing versus disordered packing267-270. Dong et al. 271

described a multidomain peptide (MDP) which utilized molecular frustration to

control the length of the self-assembled nanofibres (Figure 2.18). The peptides

are organized into an ABA block motif in which the central B block is composed

of alternating hydrophilic and hydrophobic amino acids (glutamine and leucine,

respectively) which allows the amino acid side chains to segregate on opposite

sides of the peptide backbone when it is in a fully extended β-sheet conformation.

In water, packing between two such peptides stabilizes the extended

conformation by satisfying the desire of the leucine side chains to exclude

themselves from the aqueous environment. An intermolecular β-sheet

hydrogen-bonding network can then be created between two or more pairs of

these peptides eventually growing into long fibres. Flanking Block A consists of a

variable number of positively charged lysine residues whose electrostatic

repulsion works against the desire of B block to assemble. By balancing the

forces of block A against B, it is found that at neutral pH, K2(QL)6K2 can

self-assembles into β-sheets which are soluble in water. This observation is rare

among peptides that form β-sheet assemblies which tend to generate insoluble

materials. This architectural motif may be utilized for novel tissue regeneration

strategies and other systems which require control over chemical organization at

the nanoscale.

74

Figure 2.18 (a) Primary structures of the K2 and (QL)6 series of peptides showing

the comparative domain size. (b) Proposed model of nanofibre self-assembly

indicating hydrophobic packing region, axis of hydrogen bonding, and repulsive

positive charges. 271

Another design rule that derived by copying nature are α-helices which for the

purpose of self-assembly are used as components of coiled coils272, 273. The

Wolfson group274 has demonstrated a design of an extended coiled-coil fibre

based system using sticky “end” assembly (Figure 2.19). A heterodimeric parallel

coiled coil SAF-p1 (Self-Assembling Fibre Peptide 1) and SAF-p2

(Self-Assembling Fibre Peptide 2) was designed to have a staggered hydrophobic

interface. Due to the staggered nature of the system a peptide overhang or “sticky

end” is available to form a coiled-coil interface with another peptide and so

propagate the structure along the long axis of the coiled-coil. Electron

microscopic observation revealed a much thicker fibre than expected which was

due to the lateral association between the coiled coils.

75

Figure 2.19 Computer modelling of the designed self-assembling fibre. (a) SAF-p1

(coloured yellow-to-red from the N- to the C-terminus) and SAF-p2 (coloured

blue-to-cyan from the N- to the C-terminus) interact through core residues

including asparagine pairs (coloured green) to form the two strands of a staggered,

parallel, coiled-coil fibre. (b) Negatively charged glutamate side chains (coloured

red) and positively charged lysine side chains (coloured blue) form complementary

charge interactions between SAF peptides. For clarity in this panel, both SAF

peptides are shaded dark to light gray from their N-terminus to the C-terminus. 274

2.3.2.2 Peptide amphiphiles

Amphiphilicity is one of the main driving forces for self-assembly. Molecules

containing both polar and apolar elements tend to minimize unfavourable

interactions with the aqueous environment via an aggregation process, in which

the hydrophilic domains become exposed and the hydrophobic moieties remain

shielded. Many new structures based on peptide amphiphiles (PAs) have been

synthesized with their structural and functional features explored because of the

chemical diversity which can be tolerated and their potential application in

biomimetic mineralization 275 and 3D cell culture276. Incorporation of chemical

and biochemical functionality into amphiphilic peptides has led to the

development of new surfactants with unprecedented properties leading to novel

applications in the fields of materials and biomedical research.

2.3.2.2.1 All-amino acid peptide amphiphiles

76

An important class of amphiphilic peptides is amphipathic sequences. These

sequences constitute solely of amino acids and are comprised of both

hydrophobic and hydrophilic domains when the peptide is appropriately folded. A

number of toxins like magainins277 and antibiotics such as Gramicidin S278 belong

to this class of peptide amphiphiles. To obtain a better understanding of their

assembly behaviour, research has focused on the design of all-amino acid based

amphiphilic model compounds.

Vauthey et al.279 designed a class of amphiphilic surfactant-like peptides with a

hydrophilic head group of negatively charged aspartic acid (D) at the C terminus,

thus containing two negative charges (one from the side chain carboxyl group and

the other from the C terminus) and a lipophilic tail made of hydrophobic amino

acids such as alanine (A), valine (V), or leucine (L). The N terminus is acetylated,

making it uncharged. When dissolved in water, these surfactant-like peptides tend

to self-assemble to isolate the hydrophobic tail from contact with water. Similar

to lipids and fatty acids, the supramolecular structure is characterized by the

formation of a polar interface that sequesters the hydrophobic tail from water.

(Figure 2.20)

77

Figure 2.20 Potential pathway of V6D peptide nanotube formation. Each peptide

monomer is 2 nm, and the diameter of the modeled bilayer nanotube is 50 nm. Red,

hydrophilic head; blue, hydrophobic tail. Each peptide may interact with one

another to form closed rings, which in turn stack on top of one another ultimately

yielding a nanotube. Three nanotubes are connected to each other through a

three-way junction.279

Another class of self assembling peptide is comprised of alternating hydrophobic

and hydrophilic moieties. Zhang et al.280 have described previously a class of

ionic self-complementary peptide with the sequence

(Ala-Glu-Ala-Glu-Ala-Lys-Ala-Lys)2 that spontaneously assembled to form

macroscopic membrane upon the addition of salt. The alternating hydrophilic and

hydrophobic sequence has a tendency to form an unusually stable β sheet

structure in water. When the peptides form a β sheet, they exhibit two surfaces, a

78

hydrophilic surface consisting of charged ionic side chains and a hydrophobic

surface with hydrophobic side chains. As a result, the self-assembly of these

peptides is facilitated by electrostatic interactions on one side and the

hydrophobic interaction on the other, in addition to the conventional β sheet

hydrogen bond along the backbones.

2.3.2.2.2 Lipidated peptides

The second category of peptide amphiphiles is constituted by hydrophilic peptide

sequences coupled to hydrophobic alkyl chains. It has been found that for many

single tail PAs the self-assembly leads to the formation of nanofibres which are

structurally similar to cylindrical micelles, in which the hydrophobic tails pack in

the core of the fibre while the hydrophilic peptide is displayed on the fibre’s

surface.

Since the discovery of these fibre forming PAs, their self-assembly process was

thought to occur mainly as a result of the hydrophobic interactions between

aliphatic carbon tails281. However, the work done by Paramonov et al.282

highlighted the importance of hydrogen bonding in addition to hydrophobic

packing for the stability of the nanofibres (Figure 2.21). They employed a series

of N-methylated PA molecules and it was found that blocking hydrogen bonding

in the first four amino acids after the alkyl tail prevented the gel formation and

resulted in the spherical micelles. While preventing hydrogen bonding of the

outer amino acids appeared to have little effect on the formation of a gel and thus

the molecular structure. The CD and FT-IR results suggested that the four amino

acids located in the interior of a nanofibre formed hydrogen bonds which

resemble parallel β-sheet-type interactions, while the amino acids situated on the

outer regions was able to adopt a variety of configurations.

79

Figure 2.21 (A) chemical structure of a PA which includes three distinct regions: a

hydrophobic alkyl tail, a glycine containing region, and a charged head group. (B)

Three-dimensional representation of the regions within the PA nanofibre. Region (a)

is the hydrophobic core composed of aliphatic tails. Region (b) is the critical β-sheet

hydrogen bonding portion of the peptide. Region (c) is the peripheral peptide region

which is not constrained to a particular hydrogen bonding motif and forms the

interface with the environment. 282

2.3.2.3 Aromatic short peptide derivatives

In the systems described above, the peptide chains usually contain at least ten

amino acids. However, much smaller peptides can be used by attaching aromatic

components to the peptides to take advantage of π-stacking interactions283-287.

Reches and Gazit have previously described the self-assembly of the aromatic

dipeptides FF (Phe-Phe) into straight nanotubes through π stacking288. Xu and

coworkers 289-291 were the first to report that certain Fmoc protected amino acid

and dipeptides can also spontaneously form fibrous scaffolds. The self-assembly

was triggered by enzymatic dephosphorolation of Fmoc-Yp which does not

self-assemble due to electrostatic repulsion by the charged phosphate groups. The

spectroscopic studies provide evidence for the possible molecular arrangements

in the hydrogel. The fluorenyl group stacked with both phenyl group and

themselves (in both antiparallel and parallel fashion) through π-π interactions

(Figure 2.22), which facilitate the formation of supramolecular polymers.

Self-assembly

A

B

80

Figure 2.22 Some of the possible modes of π-π interactions that contribute to the

emissions in the gel phase. 289

Recently Jayawarna et al.276 proposed a novel molecular architecture based on the

self assembly of Fmoc-FF. Smith et al.292 constructed a model of the

supramolecular structure based on a number of spectroscopic characterizations.

The proposed model accounts for the observed β-sheet signals for the peptide

portion of the molecule and the fluorescence signal from the aromatic portion of

the molecule. According to this model, the peptides are arranged as an

anti-parallel arrangement of β-sheets with the Fmoc groups acting like a zipper to

bring neighbouring sheets together, and because of the twist introduced by the

β-sheet structure the neighbouring sheets are rotated in relation to one another

creating a cylindrical structure (Figure 2.23A). These cylinders then line up

side-by-side to form a flat ribbon as observed through TEM. (Figure 2.23B) This

example and other closely related aromatic short peptide derivatives that are

known to form fibrous hydrogels have found applications in biological sensing 283

and 3D cell culture286.

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Figure 2.23 (A) A model structure was created of Fmoc-FF peptides arranged into

an anti-parallel β-sheet pattern (i) which then come together through π–π

interactions between the Fmoc groups (in orange) (ii) like a zipper to create a

cylindrical structure (iii & iv) (B) TEM image of the Fmoc-FF hydrogels composed

of flat ribbons made up of side-by-side packing of the fibrils. 292

2.3.3 Controlled self-assembly of peptides

A major challenge in molecular self-assembly is to control the self-assembly

process. Many single or multi-component systems suffer from the problem that as

soon as the material is placed into water or buffer, self-assembly is initiated.

Therefore, there is major interest in developing systems that assemble on demand.

So far, several ways have been applied to trigger self-assembly including changes

in ionic strength, pH, and temperature, light stimulus, and addition of certain

chemical entities.291, 293-301

2.3.3.1 PH/ionic strength triggered

Application of pH switch is possibly the simplest way of inducing peptide self

assembly through the elimination or reduction of the net charges presented on the

peptide molecules. A number of designed peptides are inherently pH sensitive

due to the incorporation of amino acids with ionic side chains and thus their self

assembly behaviour is, in part, governed by the pH in relation to the pKa values

A B

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of the amino acid residues. Besides, raising the ionic strength of a solution will

also achieve the same outcomes via shielding the electrostatic forces, therefore

removing pH responsiveness.302

2.3.3.2 Enzyme triggered

Enzymes have recently emerged as tools to control peptide self assembly under

physiological conditions by converting non-assembling precursors into

self-assembling building blocks under constant conditions of pH, ionic strength,

temperature. This has lead to their application in biomedical context.

Enzyme-assisted self assembly can be achieved either by catalysing the synthesis

of a self-assembly molecule303, or by removing a blocking group from a molecule

to allow assembly.

Xu and co-workers 289 reported the first example of an enzymatically driven

self-assembly process. They used alkaline phosphatase to remove a charged,

hydrophilic phosphate group from Fmoc-tyrosine phosphate to induce

hydrogelation. Based on this work, they further described the use of a

phosphatase to trigger the self-assembly of a derivative of β-amino acid, resulting

in the formation of supramolecular hydrogels, which exhibit excellent in vivo

biostability.304 Das et al.305 demonstrated two complementary enzymatic

approaches to prepare peptide nanomaterials based on aromatic short-peptide

derivatives: i) subtilisin-driven self-assembly via hydrolysis of Fmoc-peptide

methyl esters (Figure 2.24) and ii) thermolysin-driven self-assembly via amide

bond formation (reversed hydrolysis) of Fmoc-peptide-esters (Figure 2.25). The

morphology and dimensions of the nanostructures depend significantly on both

the route of self-assembly and the chemical nature of the building blocks.

83

Figure 2.24 (A) Suspension of Fmoc-Leu2-OMe and inversion of glass vial

demonstrates self-supporting gel formation of Fmoc-Leu2 after ester hydrolysis

using subtilisin (Entry 1). (B) Proposed mechanism of Fmoc-peptide ester

hydrolysis that self-assembles to form higher-order aggregates through π–π

interlocked β-sheets. 305

Figure 2.25 Solutions of Fmoc-Thr-OH and Leu-OMe. The inversion of the glass

vial demonstrates self-supporting gel formation of Fmoc-Thr-Leu-OMe via

reversed hydrolysis by thermolysin (entry 6). 305

A reversible system using two enzymes, kinase and phosphatase that catalyse

opposite reactions of phosphorylation and dephosphorylation respectively, was

demonstrated by Xu and coworkers.306 The peptide hydrogelator used is a

naphthalene linked to FFGEY (Nap-FFGEY, Figure 2.26i) which forms gel in

water (Figure 2.26iii (A)). Adding a kinase to the hydrogel induces a gel-sol

phase transition (Figure 2.26iii (B)) in the presence of adenosine triphosphates

(B)

(A)

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(ATP) since the tyrosine residue is converted into tyrosine phosphate and further

treating the resulting solution with a phosphatase removes the phosphate group

(Figure 2.26ii) and restores the hydrogel (Figure 2.26iii (C)). This biomimetic

approach promises a new way to regulate the state and properties of

supramolecular hydrogels.

Figure 2.26 (i) Chemical structure of Nap-FFGEY. (ii) Reversible modification of

the peptide gelator by a phosphatase/kinase reaction. (iii) Optical images of (A) gel

formed initially (B) the solution obtained after adding a kinase to A (C) gel restored

after adding a phosphatase to B. 306

2.3.4 Mineralization

2.3.4.1 Biomineralization Biomineralization is the process by which living organisms produce minerals,

such as those found in mollusk shells,307,308 sea urchin spine,309,310 teeth311 and

diatom cell walls312. Natural mineralization creates the most intricately stunning

inorganic structures and in marked contrast to anthropogenic and geological

syntheses of minerals which typically require elevated temperatures or extremes

of pH, living systems produce such exquisite structures under mild physiological

conditions of ambient temperatures and near-neutral pH. Bone is a particular

complex example in which collagen molecules serve as nucleation sites for

(i) (ii)

(iii)

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oriented crystals of hydroxyapatite (HA). The collagen fibrils are formed by

self-assembly of collagen triple helices and the HA crystals grow within these

fibrils in such a way that their c axes are oriented along the long axes of the fibrils

resulting in the excellent mechanical properties of bones.313

Such a high degree of organization is achieved, in part, through the presence of

organic macromolecules, which control many aspects of the mineralization

process, including crystal nucleation and growth.314 Therefore, understanding the

design of the biological macromolecules is vital to achieve the same outcome in

vitro.

Recent investigation of the mechanisms governing the biological synthesis of

silica in a marine sponge Tethya aurantia led to the surprising discovery that this

process is catalysed by a protein called silicateins (for silica proteins), which are

the subunits constituting the central protein filament (1–2 mm in length and 2 um

in diameter). These axial filaments of protein are occluded within the silica

spicules (1–2mm in length and 30 um in diameter) which account for 75% of the

dry weight of the organism. Characterization of silicatein α (the principal subunit

comprising nearly 70% of the mass of the filaments) and its cloned DNA

indicated that it is highly homologous to members of the cathepsin L subfamily of

the papain family of proteases, including similarities between their amino acid

sequences and three-dimensional structures. Besides, of the three residues of the

“catalytic triad” of the cathepsin active site, two, His (imidazole side chain) and

Asn, are conserved in silicatein α, but the third active-site residue in cathepsin,

Cys (sulfhydryl side chain), is replaced in silicatein α by Ser-26 (hydroxyl side

chain). At this position, the structure of silicatein α resembles that of the serine

proteases315.

Cha et al. 316 have demonstrated that the silicatein filaments and their constituent

subunits that were isolated from the sponge could catalyze the hydrolysis and

direct the polymerization of silica and silicone polymer networks both chemically

and spatially in vitro, under conditions at ambient temperatures, pressure and

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neutral pH. Based on the structural homology between silicateins and the

corresponding proteases and that both the condensation of silicon alkoxides

promoted by the silicateins and the cleavage of peptides catalyzed by the

proteases must proceed through an obligatory hydrolysis reaction, and that both

are known to be accelerated by general acid–base catalysis, the author suggested

that silicatein α catalyzes the siloxane polymerization reaction through the

activity of the serine and histidine side chains that occupy positions

corresponding to the catalytically active, functionally related side chains in the

proteolytic enzymes of both the cathepsin L (cysteine - histidine) and

trypsin/chymotrypsin (serine-histidine) types. This was further supported by the

subsequent site-directed mutagenesis studies which confirmed the requirement for

both Ser-26 and His-165 residues to present at the active site of silicatein α for

efficient catalysis of alkoxysilane polycondensation at neutral pH. Besides, the

catalytic efficiency is also dependent on the precise three-dimensional

conformation of the native protein 317. Such a mechanism is illustrated in Figure

2.27 This mechanism may be harnessed for the development of environmentally

benign routes to the synthesis of patterned silicon-based materials.

87

Figure 2.27 Proposed mechanism of silicon ethoxide condensation catalyzed by

silicatein α. Hydrogen bonding between the imidazole nitrogen of the conserved

histidine and the hydroxyl side chain of the active-site serine is proposed to enhance

the nucleophilicity of the serine oxygen atom, potentiating its attack on the silicon

atom of the substrate. Nucleophilic attack on the silicon atom displaces ethanol,

forming a transitory protein-O-Si intermediate that potentially would be stabilized

as a pentavalent silicon species through a donor bond from the imidazole nitrogen

atom. Upon addition of water, the intermediate is subjected to hydrolysis, resulting

in the generatin of the reactive silanol, and restoration of the serine-histidine

hydrogen-bonded pair at the enzyme’s active site. Condensation initiated by

88

nucleophilic attack of the released Si-O- on the silicon of the second substrate

molecule then forms the disiloxane product. 316

Hecky et al. 318 have studied the amino acid composition of the diatom cell-wall

protein which serves as the template for silica deposition. It was found to be

enriched in serine plus threonine (hydroxyl-containing amino acids), suggesting

that the hydroxyl group may mediate the silicification process in the diatom

cell-wall. The function proposed for the template protein is presented

schematically in Figure 2.28. The protein template presents a layer of hydroxyl

groups which can undergo condensation reactions with silicic acid molecules.

The initial layer of condensed silicic acid will then be held fixed to the protein

template, in a geometric arrangement that will favour further polymerization of

silicic acid to form the silica-rich frustule. This is kinetically more favorable than

simply allowing the silicic acid molecules to come together by random collisions.

This proposed mechanism was further supported by the thermodynamic

calculations 319. Such a mechanism may also contribute to the biosilicification in

marine sponge.

Figure 2.28 Proposed condensation reaction between silicic acid and serine on the

protein template of the silicalemma. Water by-product may be eliminated or

structurally incorporated into the forming frustule through hydrogen bonding with

the oxygens of silica. 318

However, the simple density of hydroxyls alone is not sufficient for

polymerization of the silicon alkoxides as evidenced by the lack of activity of the

hydroxyl-rich cellulose and silk polymers. The conformation of such groups in

the silicatein molecule may also play an important role.

2.3.4.2 Biomimetic mineralization

89

The success of nature in creating such well-defined and precisely controlled

inorganic structures is a source of inspiration for creating synthetic counterparts,

especially at the nanoscale. Furthermore, there is a growing demand for benign

synthetic conditions that will minimize adverse environmental effects.

Inspired by these natural examples of silicification, several synthetic templates

including polyamines 320-323, polypeptides324-328, peptide-polymer hybrids329 and

self-assembled peptides330-332 have been demonstrated for mimicking the catalytic

activity of silicateins and silaffins under ambient conditions, yielding hybrid

materials with diverse morphologies depending on factors such as temperature

and pH of the reaction medium333. Cha et al.334 have reported the use of

self-assembled cysteine-lysine block copolypeptides as templates for the

production of mesoporous silica spheres and tightly packed silica columns. The

nucleophilic group of cysteine is shown to be essential for the hydrolysis of the

silicon alkoxide, while the protonated amine group of lysine promotes the

polycondensation and deposition of silica. Systematic substitution of residues

used in the construction of these peptides showed that the rate of catalysis was

proportional to the strength of nucleophilicity of the nucleophilic side chain.

Similarly, Yuwono and Hartgerink335 have demonstrated templating of silica by

cationic peptide amphiphiles, which are β-sheet forming peptides with alkyl chain

conjugated at the N-terminus.

Based on the finding of the conserved residues in the active site of silicatein, Roth

et al.336 have evaluated an array of bifunctional small molecules containing both

nucleophilic group (such as –SH or -OH) and hydrogen-bonding acceptor group

(primary or substituted amine) for their catalytic activity in the in vitro formation

of silica from silicon alkoxides at neutral pH and room temperature. Among the

tested molecules, cysteamine and ethanolamine show the highest catalytic activity

while lacking either one of the two functionalities leads to significantly reduced

catalytic activity. The resulting silica nanoparticles showed no higher ordered

structures indicating that the small molecules have little ability to attract and align

90

nuclei, but rather act as a catalyst for the hydrolysis of silica precursors. The

potential of the small molecule catalyzed sol-gel synthesis of silica in the area of

biomaterial encapsulation was further demonstrated for enzymes, fluorescent

proteins, and live cells. The benign reaction conditions help to preserve the

activities of these encapsulated biological materials. This study support that the

two functional groups in close proximity to each other are required for catalyzing

the hydrolysis of TEOS.

Morse et al. 334,337 have demonstrated that the hydroxyl or imidazole group alone

was not sufficient to catalyze the hydrolysis of TEOS, but rather capable of

catalyzing silicic acid condensation 326, 327. However, another study using TMOS

as silica precursor have observed silica formation in the presence of

poly-histidine. 327

The translation of the mechanism to a more robust synthetic system was further

developed using golden nanoparticles (GNPs) functionalized with the appropriate

nucleophilic (hydroxyl) and hydrogen bonding (imidazole) functionalities

respectively (Figure 2.29 (A), (B), (C))337. The combination of the two

populations was found to successfully mimic the catalytic activity of silicateins.

Significant catalytic activity was also observed when imidazole functionality was

replaced with a primary amine as the hydrogen-bonding acceptor. Replacement of

either of these functionalities by a non-reactive methyl group abolished the

catalysis in this synthetic system. TEM image of the silica product showed that

the GNPs were entrapped in the silica network and the SAED pattern indicated

the amorphous nature of the silica product (Figure 2.29D).

91

Figure 2.29 Schematic of the interaction between two GNPs (B,C) capped with

imidazole and hydroxyl functionalities (A). (D) TEM image of silica product with

entrapped GNPs. Selected area electron diffraction (inset) indicating amorphous

nature of silica. 337

2.4 Graphene and graphene-based nanocomposites 2.4.1 Introduction to graphene 2.4.1.1 Structure and properties of graphene

Graphene, as the third member of the nanocarbons family following fullerene and

D

92

CNTs, can be regarded as the basic constituent of graphitic systems (Figure 2.30) 338.

It consists of a single atom layers of carbon arranged in a hexagonal lattice. In the

graphene layers each carbon atom is trigonally bonded to the three nearest

neighbors by means of sp2 hybridized orbitals which form the strong covalent

C-C (σ) bonds. The overlap of the un-hybridized p-orbitals from each carbon

atom forms the delocalized π-bonds. The overlapping orbitals produce a fully

conjugated system, providing fast electron transfer within the layers.

Figure 2.30 Mother of all graphitic forms. Graphene is a 2D building material for

carbon materials of other dimensionalities. It can be wrapped up into 0D buckyballs

(left), rolled into 1D nanotubes (middle) or stacked into 3D graphite (right). 338

This 2D analogue of CNTs has recently received tremendous interest owing to its

exceptional properties 339. The electrons in high quality graphene move both

relativistically and ballistically with mobilities exceeding ~15,000 cm2 V-1 s-1 at

room temperature, which is faster by two orders of magnitude than that of Si 340

and is about an order of magnitude slower than that of CNTs. Graphene’s

outstanding mechanical stiffness (~1060 GPa) 341 and thermal conductivity

Graphene

93

(~3000 W m-1 K-1) 342 compete with the in-plane values for graphite. Their

biocompatibility 343 and extremely high specific surface area (up to 2600 m2/g) 344

which is larger than that of CNTs (> 1000 m2/g) makes it an attractive material

for a wide range of applications including next-generation nanoelectronic

devices345,346, photonic devices 347, catalyst supports 348, composite materials349,350,

drug delivery materials 351 and sensors 352,353. Additionally, the unique

combination of their optical transparency and excellent electricity and heat

conductivity makes them ideal for thin and flexible flatscreens and suitable for

replacing the ITO in photovoltaic cells.

Although graphene exhibit no advantages over CNTs, as for applications, there

are exciting fundamental effects which can be studied in graphene and may help

to understand the nanostructures in general and indirectly advance

nanotechnology. Furthermore, several studies have demonstrated that at low

nanofiller content graphene perform significantly better than CNTs in terms of

enhancing the mechanical properties of the nanocomposites 354. The superiority of

graphene platelets over CNTs in terms of mechanical properties enhancement is

mainly attributed to three distinct advantages of graphene: (1) their larger specific

surface area for considerably more contact with the matrix material than the

tube-shaped CNTs; (2) improved nanofiller-matrix adhesion benefiting from the

mechanical interlock of the wrinkled surface of graphene with the surrounding

matrix material; (3) the 2D geometry of graphene sheets are shown to be more

effective in suppressing crack propagation in polymers than the 1D

nanotubes354,355.

Du et al. 356 have demonstrated that the improving effect of graphene nanosheets

(GNSs) as conducting fillers on the electrical conductivity of their composites

was far lower than theoretically expected 357. GNS/high density polyethylene

(HDPE) composites showed much higher percolation threshold and lower

electrical conductivity than the MWNT/HDPE composites at the same filler

content. Regardless of the similar conductivity of the two fillers, these results can

94

be attributed to several reasons: (1) MWNTs were much easier to be isolatedly

dispersed within the HDPE matrix than GNSs; (2) the 2D structure of GNSs was

hardly to be maintained in the preparation of the composites; (3) the 2D GNSs

were not as effective as the 1D MWNTs in forming the conductive networks.

Therefore, to make a full use of the advantages of the 2D GNSs as conductive

fillers, measures such as the development of more effective composite

preparation techniques to avoid GNS rolling and aggregation and the construction

of the desired contact among GNSs are highly demanded.

While monolayer graphene attracts the most attention since their unusual

electronic band structure allows its carriers to behave as massless Dirac fermions,

there is also significant interest in few-layer graphene. Owing to their distinct

electronic band structures, AB-stacked bilayer 358 and trilayer 359 graphene have

extraordinary potential for next-generation optoelectronic and microprocessor

applications. However, the lack of bandgap for graphene significantly limits their

potential in electronics applications. Recently, great progress have been made on

the creation and precise tuning of a bandgap in bilayer graphene by applying a

vertical electric field 360-364 which introduce asymmetry into the bilayer structure,

thus leading to the formation of a bandgap. Trilayer graphene with ABC

crystallographic stacking was also reported to exhibit an induced bandgap under

the application of a perpendicular electric field 365. Single-layer graphene is

intrinsically semimetal, introducing an energy band gap requires patterning

nanometer-width graphene ribbons 366 or utilizing special substrates 367,368.

Furthermore, compared with monolayer, bilayer graphene has a lower electrical

noise and an intrinsic screening of the influence of trapped charges in the gate

dielectric. Such breakthrough opens up new possibility of bi and trilayer graphene

based electronics and photonics.

2.4.1.2 Production of graphene

Ever since its isolation in 2004 by mechanical cleavage of graphite 339, a wide

range of techniques have been reported for the synthesis of graphene. In general,

95

these techniques can be divided into ‘bottom-up’ and ‘top-down’ methods. The

bottom-up methods usually include epitaxial growth of graphene films on

single-crystal SiC 369,370 and chemical vapor deposition (CVD) on catalytic

metals371-373. These methods are not widely used because of their complexity,

limited scaling-up and high cost of the specialized fabrication systems. Large

scale production of high-quality graphene is usually achieved by the top-down

methods including liquid-phase exfoliation and thermal expansion of graphite 374.

Among these techniques, liquid phase exfoliation which includes the reduction of

exfoliated GO and sonication-assisted direct exfoliation of graphite in solution

are particularly attractive owing to their low cost and massive scalability 375,376.

They also facilitate the processing of graphene as well as offer the flexibility to

deposit them onto any desirable substrates. In addition, the reduction of GO

method benefits the further modification of graphene surface for the development

of functionalized graphene-based materials that hold huge potentials in energy

storage 377, catalysis, biosensing 378 and drug delivery 379. The advantages and

disadvantages of each of the methods were briefly reviewed below.

2.4.1.2.1 Micromechanical cleavage

Fabrication of graphene from graphite requires exfoliation of the individual layers.

This was originally achieved by micromechanical exfoliation, simply using

‘scotch tape’ to peel off micrometer-wide graphene sheet from HOPG 339.

Micromechanical cleavage produces the highest quality of graphene which are

suitable for fundamental studies. However, the extremely low productivity owing

to the manual effort combined with its high cost and failing to scale up for

commercially viable devices has considerably limit their practical applications.

2.4.1.2.2 Liquid phase exfoliation

Liquid phase exfoliation has been extensively developed in recent years for mass

production of graphene. Among the solution phase strategies, the most popular

route is the reduction of exfoliated GO to produce graphene-like nanosheets best

known as reduced graphene oxide (RGO) 375, 376. Three steps are commonly

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involved (Figure 2.31): (1) Oxidation of graphite to hydrophilic graphite oxide in

which graphene basal planes are decorated with covalently bonded oxygen

functional groups. (2) The oxygenated graphite is exfoliated in aqueous solution

to give monolayer graphene oxide (GO). (3) Subsequently, GO nanosheets are

reduced chemically 375 or thermally 380 to partially restore the sp2- hybridized

network 381 to yield RGO or chemically converted graphene (CCG) 382.

The three routes that are typically employed for obtaining graphite oxide are the

Brodie383, Hummers’ 384 and Staudenmaier methods 385. The Hummers’ method is

the most commonly used due to its relatively short reaction time and absence of

hazardous gas during the synthesis. The GO suspension is indefinitely stable

owing to the polar nature of the functional groups introduced onto the graphene

sheets. Their good water dispersibility makes it easy to process them in aqueous

and polar solvents thus offers the convenience to cast thin films for various

technological advances 386. The presence of these oxygen-containing functional

groups not only renders GO good water solubility but also provides reactive sites

for the deposition of metal 387 and inorganic nanoparticles. GO is also the starting

material for most other chemically modified graphenes (CMGs), which are of

interest in their own right 388-390. GO is favorable for its tunable electronic and

chemical properties. Different degrees of oxidation 382 and species of oxygen

containing groups 391 would lead to a diverse energy gap and structure

distortion392, and thus cause different conductive and chemical properties of

graphene. TEM studies on GO revealed that the underlying carbon lattice

maintains the order and lattice spacings of graphene. It also shows that single GO

sheets are highly electron transparent and stable in the electron beam and

therefore ideal for use as TEM support films393.

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Figure 2.31 Preparation of graphene by chemical reduction of GO synthesized by

Hummers’ method.

Although there is still much debate over the structure of GO (with uncertainty in

both the type and distribution of the oxygen-containing functional groups), the

most prevalent model currently appear to be 1, 2 epoxides and hydroxyls

distributed across the basal plane, with carboxyl groups at the edges

(Lerf–Klinowski model, Figure 2.32) 388. Studies have also demonstrated that the

fully oxidized GO has an atomic C/O ratio of roughly 2:1388.

(GO)

98

Figure 2.32 Schematic model of a GO sheet, with -COOH hanging on the edge and

-O- and –OH decorate the basal plane. 388

Recently, Wilson et al. 394 have proposed a new structure of GO by analogy with

oxidized CNTs. As shown in Figure 2.33, the GO comprises of oxidative debris

(OD) which non-covalently decorates oxidized graphene sheets (denoted as

base-washed GO, bwGO). These oxidative debris act as a surfactant to stabilize

the aqueous suspension and can be removed by washing with base. The OD was

found to be made up exclusively of small, low molecular weight compounds and

formed at least 30% of the weight of the original “graphene oxide”. Although the

degree of oxidative functionality is clearly reduced in bwGO compared with

as-produced GO, it is still significant. The electrical conductivity of bwGO film is

shown to be very similar to that of reduced GO.

Figure 2.33 Schematic representation of as-produced GO: large oxidatively

functionalized graphene-like sheets with surface-bound debris. Note that the

graphene-like sheets extend further than depicted. 394

99

The reduction of GO provides an efficient route for graphene production in terms

of commercial scalability and high throughput of predominantly monolayers with

large flake size, opening the door to integrate graphene with other materials to

form nanocomposites with improved performance 395,396. However, they still face

some drawbacks. For example, due to the incomplete removal of oxides397 as well

as the structural defects induced 375,382, the electronic property of RGO is still

largely degraded, which is several orders of magnitude lower than that of pristine

graphene 382,385, making them unattractive for electronics applications. Thus, a

non-covalent method for producing significant quantities of defect-free,

unoxidized graphene is urgently required.

The other solution phase pathway employs ultrasonic energy to directly exfoliate

graphite down to thin flakes 398-403. Exfoliation in both organic solvents and

aqueous surfactant solutions has been reported.

The trick to this method is to choose the solvents whose surface energy matches

well with that of graphite, thus allowing for strong interaction of the solvent with

graphite and minimizing the energy required for exfoliation. Commonly used

solvents include N-methyl-pyrrolidone (NMP) 400 and dimethylformamide

(DMF)347. However, such method tend to suffer from a major problem: the low

concentration of graphene in the solution that can be achieved 400 which is a

significant barrier to many of the applications, such as for producing

graphene/polymer composites 404. In order to maximize the concentration of

graphene while maintaining their quality, Coleman and coworkers 401 have

applied mild sonication for extended times (up to 146 h) in solvent NMP to

obtain graphene flakes of <10 layers at a concentration of up to 1 mg mL-1. The

mean flake length remains above 1 μm for up to 460 h of sonication. They further

employed a two-step sonication and successfully obtained the good-quality

dispersion at the concentration of 17 mg mL-1 which is the highest ever achieved

so far 405.

100

As the solvent–graphite interaction is van der Waals, solvent exfoliated graphene

exhibit minimal basal plane defects 400, therefore the property of pristine graphene

is largely preserved. It also provides a viable route for mass production of

graphene on a variety of substrates not available using cleavage or growth

methods, thus facilitate production of graphene based composites and films. The

defects present on RGO 406 and solvent exfoliated graphene 400 have distinguished

them from both the ideal graphene crystal and one another. While the deviations

from perfection deteriorate the performance of graphene-based devices, the

unique properties introduced owing to the presence of defects may lead to

interesting effects and potential applications 407.

However, this process is not without its drawbacks. These solvents are expensive

and require special care when handling. In addition, they tend to have high

boiling points, making it difficult to completely remove them. This can present

problems for flake deposition and composite formation. Coleman et al. have

recently reported preparation of graphene dispersion in low boiling point solvents

such as chloroform and isopropanol 408. Although using these low boiling point

solvents allows individual flakes to be deposited on substrates, the organic

solvent limit their application in bio related area. Unfortunately, the most useful

solvent of all, water, has a surface energy that is much too high to work on its

own as an exfoliant for graphene.

Surfactants mediated exfoliation of graphite has been recently developed 399,402,403.

Although compared with solvent exfoliation route, the concentration of the

dispersion and the percentage of monolayers present are low, it avoids the use of

toxic solvents and the aqueous environment brings its own advantages, such as

the ability to deposit individual flakes onto surfaces. Commonly used surfactants

for CNT dispersion, such as SDBS, have been employed for graphite exfoliation 402.

Coleman et al. have also demonstrated that for ionic surfactants, the

concentration of the dispersion is proportional to the magnitude of the

electrostatic potential barrier which stabilizes surfactant coated flakes against

101

aggregation, while for the non-ionic surfactants, the dispersed graphene

concentration scaled linearly with the magnitude of the steric potential barrier

stabilizing the flakes 409.

The Coleman group further extended the breakthrough to produce graphene

aqueous dispersion at a concentration of 0.3 mg mL-1 by using sodium cholate as

surfactant 403. Sodium cholate is believed to be a far-superior surfactant than

SDBS in terms of the concentration and the degree of exfoliation achievable. The

dispersions (with graphene concentration as high as 0.04 mg mL-1) were also used

to form transparent thin films which exhibit a direct current conductivity of up to

1.5x104 S m-1 benefiting from their oxide and defect-free natures 399.

For surfactant-graphene dispersion, the concentration above 1 mg mL-1 has not

been achieved 403. Another significant issue facing is that this method tend to

yield size- and layer-polydispersed flakes, out of which a large population of

small flakes with lateral size of ~1 μm or less are produced that were useless for

many applications. To this end, a study has demonstrated the successful isolation

of monodisperse colloidal solutions of single graphene sheets and few-layer

graphene stacks from bulk graphene dispersions via density-gradient

ultracentrifugation 410. Khan et al. have also reported a size selection method by

controlled centrifugation 411.

2.4.2 Graphene based nanocomposites and nanohybrids

Among the diverse applications of graphene, a particularly active area of research

concerns the development of graphene-based composite and hybrid materials,

which seeks to combine the fascinating properties of the two-dimensional

nanocarbon with additional functionality provided by a second component.

Various metals including Au 412, Ag 413, Pt 414 and metal oxides 415-418 have been

anchored to the surface of graphene for the applications as electrochemical

sensors in which graphene serves as a support material providing larger

102

electrochemically active surface areas.

Especially, by combining the unique properties of graphene with the

well-established photoactivity of TiO2, the development of their composites and

hybrids with synergistic effect have found great potentials in the fields of

photocatalysis, dye-sensitized solar cells and hydrogen evolution. The remarkably

enhanced photocatalytic activity of graphene-TiO2 nanocomposites compared

with pure TiO2 was attributed to the high specific surface area of graphene for

high dye adsorption capacity and their role as electron acceptor to effectively

hinder the electron–hole pair recombination of TiO2 419. Furthermore, graphene

may exhibit the same photosensitizing effect as CNTs to enhance the utilization

of visible light in photocatalysis122. Graphene-TiO2 composites also showed

higher photocatalytic efficiency over CNT-TiO2 composites benefiting from the

better conductivity and higher surface area of graphene.

Many efforts have been made to fabricate graphene-TiO2 nanocomposites and

hybrids. The challenge is to uniformly integrate graphene sheets into the matrix

materials or to achieve uniform dispersion of the loaded TiO2 nanoparticles.

Strategies such as atomic layer deposition 420 and electron beam irradiation 421

have been reported, with the latter requires pretreated graphene as a raw material.

Meng et al. 420 have employed atomic-layer-deposition method for the

preparation of graphene-TiO2 nanocomposites with precisely controlled

morphology and phase of TiO2. However, since it is a gas-solid synthesis route, it

lacks the flexibility to deposit the nanocomposites on arbitrary substrate. Besides,

the process is probably not commercially viable since it requires a six-step

sequence with heat treatment at various temperatures.

Alternatively, methods including self-assembly 422, molecular grafting 423 and

sol-gel process 424 have been employed for the production of graphene-TiO2

nanocomposites and hybrids. Wang et al. 422 reported using SDS stabilized FGSs

(functionalized graphene sheets) for in-situ growth of TiO2 nanoparticles and the

103

resulting hybrid materials showed significantly enhanced Li-ion

insertion/extraction capacity in the hybrid electrodes at high charge/discharge

rates due to its high surface area and excellent conductivity, making graphene

sheets highly promising for Li-ion battery electrode materials. Tang et al.423

incorporated chemical exfoliated graphene sheets (GS) into TiO2 nanoparticle

films via a molecular grafting method for dye-sensitized solar cells. The

composite films showed significantly enhanced conductivity and photovoltaic

performance attributed to the creation of a continuous electron transfer network

as a result of the implanted GS, which lower the probability of recombination of

photoinjected electrons and the better dye loading of GS/TiO2 film.

Chen et al. 425 reported the fabrication of graphene sheets-wrapped anatase

hollow particles by first functionalizing the electroactive TiO2 hollow particles

with aminopropyltriethoxysilane to obtain a positively-charged surface. The

negatively charged GO sheets were then linked to these functionalized TiO2

hollow particles via electrostatic interaction. Finally, the GO sheets were reduced

to graphene sheets by thermal treatment, leading to the formation of

graphene–TiO2 composites. Similarly, Kim et al. 426 have prepared the core/shell

structure of nanographene sheets self-assembled onto TiO2 nanoparticles and it

was found to show higher photocatalytic and photoelectrochemical activities than

that of the conventional composites which has TiO2 nanoparticles loaded on the

micrometer-sized graphene sheets. This is attributed to the improved contact

between r-NGO and TiO2 NPs.

Kamat et al. 427,428 have reported that GO underwent photocatalytic reduction as it

accepted electrons from UV-irradiated TiO2 in GO-TiO2 nanocomposites (Figure

2.34). The work further revealed that approximately 50% of the oxygen sites are

able to accept electrons from TiO2 and undergo reduction. Thus the photocatalytic

reduction of GO-TiO2 opens up new ways to obtain graphene-semiconductor

composites that meet the requirement to both obtain graphene as individual sheets

and to maintain it in the reduced form. The advantage of photocatalytic reduction

104

over the conventional chemical reduction is that (1) use no toxic chemicals such

as hydrazine (2) it can be triggered on demand by tuning the UV-irradiation and

(3) the binding of the oxide particles keeps the exfoliated graphene sheets from

collapse after reduction while the chemical reduction by hydrazine induce the

irreversible reaggregation of graphene sheets 375.

Figure 2.34 TiO2-graphene composite and its response under UV-excitation. 427

The pioneering study by Kamat has stimulated an extensive range of studies on

the fabrication and application of graphene-TiO2 photocatalysts from the

reduction of GO-TiO2 nanocomposites. Zhang et al. 429 have reported the

preparation of a chemically bonded P25-graphene nanocomposite using GO and

P25 as starting reactants via a one-step hydrothermal method, during which both

of the reduction of GO and loading of P25 were achieved. The resulting

composite exhibited a significant enhancement of photocatalytic degradation of

methylene blue (MB) under both UV and visible light irradiation over bare P25.

Furthermore, G-P25 showed higher photocatalytic efficiency than P25-CNTs

with the same carbon content mainly due to its giant two-dimensional planar

structure, which facilitated a better platform for adsorption of dyes and charge

transportation.

Zhou et al. 430 reported the one-pot in-situ preparation of graphene-TiO2

nanohybrids via solvothermal reaction using GO and TBOT as starting materials.

The reduction of GO to graphene in GO-TiO2 nanohybrids via the solvothermal

105

reaction is claimed to be more effective in lowering the oxygen and defect levels

in graphene compared with using reductants like hydrazine and NaBH4 431.

Akhavan 432 reported the substantially improved antibacterial activity of the

reduced GO-TiO2 which is achieved by photocatalytic reduction of GO-TiO2

under UV-Vis light irradiation compared with that of GO-TiO2. This may due to

the fact that the reduced GO has better conductivity and thus result in better

charge separation between photoexcited electrons and holes leading to a reduction

in recombination rate of the pairs.

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127

Chapter 3 Experimental methods

3.1 Materials Unless otherwise stated, all the chemicals were purchased from Sigma-Aldrich,

are laboratory reagent grade and were used as received. Tetrakis (2-hydroxyethyl)

orthosilicate was purchased from Xuzhou Xuri Chemicals Co., Ltd. Graphite

powder (grade 2369) was purchased from Branwell graphite Ltd, UK. Alkaline

phosphatase from bovine intestinal mucosa was in the form of a lyophilized

powder at ≥2,000 DEA units/mg protein. GO and bwGO were provided by

Cristina Vallés and were synthesized via the Hummers method 1. bwGO was

prepared according to a method reported by Wilson et al. 2 . The solvents used in

the current study were supplied by Fisher Scientific. The dH2O was produced in

Materials Science Center, University of Manchester. The phosphate buffer

solution was prepared by dissolving K2HPO4 and MgCl2 in double distilled water

at a concentration of 0.6 M.

3.2 Experimental procedure

3.2.1 Synthesis of aligned CNT arrays by injection CVD

method

The aligned CNT arrays were synthesized in-house using a method previously

reported by Singh et al. 3. Briefly, the arrays were grown on a substrate of

oxidized silicon in a two-stage tube furnace using an injection chemical vapour

deposition (CVD) method. Silicon wafers (5 mm x 5 mm, N<100> Si wafers with

resistivity of 1-10 Ω⋅cm, single side polish, IDB Technologies Ltd) were

pre-cleaned successively in acetone and H2O by ultrasonication to remove surface

contaminants and leave to dry in air. The cleaned wafers were placed into a quartz

tube in the furnace. The wafers were oxidized in the heating zone of the furnace

at 760 ºC under air flow for 1 h, prior to a purge with 400 mL/min argon for 20

min to remove the remaining air in the tube. A solution of 5 wt% ferrocene

(catalyst precursor) in toluene (carbon feedstock) was injected at a rate of

128

0.04mL/min into the pre-heating stage (200 ºC) of the furnace where the solution

was vaporized and carried by argon flowing at 100 mL/min into the heating zone

of the furnace. Aligned CNT arrays were grown from the oxidized silicon

substrate for 1 h at 760 ºC. The furnace was then cooled to room temperature

under the protection of argon to avoid oxidation of CNTs. The diagram of the

furnace set-up is displayed in Figure 3.1. The morphology and dimensions of the

as-grown CNT arrays were characterized using SEM and TEM.

Figure 3.1 Schematic diagram showing the set-up for the CVD synthesis of aligned

CNT arrays.

3.2.2 Adsorption study of the surfactants on CNTs

3.2.2.1 Adsorption of the surfactants on aligned CNT arrays

The adsorption behaviour of the surfactants (which acted as surface modifiers in

the inorganic coating experiments) was studied by monitoring the change in the

concentration of the surfactant solution in the presence of the nanotubes over time

using UV-Vis spectroscopy (Hitachi U-1800 spectrophotometer, U.K). A series

of aromatic Fmoc-AAs were studied to investigate the effect of amino acid side

chain on the adsorption behaviour. Figure 3.2a shows the four aromatic

Fmoc-AAs studied: Fmoc-tryptophan (Trp), Fmoc-phenylalanine (Phe),

Fmoc-tyrosine (Tyr), and Fmoc-histidine (His). Fmoc-Glycine (Gly) was also

used as a nonaromatic control.

Argon

Syringe containing 5wt%ferrocene in toluene

Preheating stage 200 ºC

Temperature control

Silicon oil

Quartz tube CNT arrays grown on silica

substrate

Heating zone 760 ºC

129

In a typical measurement, aligned CNT arrays on a 5 mm x 5 mm silicon wafer

(Figure 3.2c) were placed into a quartz cuvette containing 3 mL of 0.05 mM

Fmoc-Trp aqueous solution (Figure 3.2b). UV-Vis measurements were carried

out over time to monitor the change in the absorbance of the solution. The spectra

were also collected in the absence of the substrate as a reference. The same

measurement was performed for all the modifiers studied. The Fmoc group was

also compared with other aromatic ligands such as the benzyl group of benzyl

alcohol (Figure 3.2a).

130

Detector

UV/Vis Light Source

Modifier solution

(d)

10μm

(c)

Fmoc group R1 = any amino acid

Histidine (His) Glycine (Gly)

Tryptophan (Trp) Phenylalanine (Phe) Tyrosine (Tyr)

R1=

(a)

(b)

Benzyl

alcohol

131

Figure 3.2 (a) Molecular structures of the modifiers studied. (b) Scheme illustrating

the UV-Vis measurement of the adsorption of the surfactant on (c) aligned CNT

arrays (side-view) and (d) randomly aligned CNT networks.

The concentration of the solution was determined by using the Beer-Lambert law.

Which describes the linear relationship between the absorbance and concentration

of an absorbing species which is usually written as:

A= ε b c (3.1)

where A is the measured absorbance of the modifier solution (dimensionless), ε is

the molar absorptivity of the modifier molecule in solution (L mol-1cm-1), b is the

path length of the cuvette, which is 1 cm in the current study, and c is the

concentration of the modifiers in solution (mol L-1). According to Beer-Lambert

law, the absorbance is directly proportional to the concentration of the solution.

However, it should be noted that the Law is not obeyed at relatively high

concentrations. A calibration curve was constructed for calculating the

concentration of the solution from the absorbance. The calibration curve was

plotted by preparing a series of modifier solutions of known concentrations which

were spaced relatively equally apart and the absorbance at 265 nm (corresponding

to the absorbance of the Fmoc group) was measured. In the case of BA, the

absorbance at 257 nm was measured. The calibration curves of all the modifiers

studied are shown in Figure 3.3.

(a)

0

0.1

0.2

0.3

0.4

0.5

0 0.005 0.01 0.015 0.02 0.025

Concentration / mM

Ab

so

rban

ce

132

0

0.2

0.4

0.6

0.8

1

0 0.005 0.01 0.015 0.02 0.025 0.03 0.035 0.04 0.045

Concentration / mM

Ab

so

rban

ce

(b)

0

0.5

1

1.5

2

0 0.02 0.04 0.06 0.08 0.1 0.12

Concentration / mM

Ab

so

rban

ce

(c)

0

0.2

0.4

0.6

0.8

1

0 0.01 0.02 0.03 0.04 0.05 0.06

Concentration / mM

Ab

sorb

ance

(d)

133

Figure 3.3 Calibration curves of all the modifiers studied. (a) Fmoc-Trp. Equation

of line: y=20.07x+0.0081, correlation coefficient R2=0.99. (b) Fmoc-Tyr. Equation of

line: y=19.01x-0.0080, R2=1.00. (c) Fmoc-His. Equation of line: y =16.69x + 0.0173,

R2=1.00. (d) Fmoc-Phe. Equation of line: y=18.49x+0.0005, R2=1.00. (e) Fmoc-Gly.

Equation of line: y=17.30x+0.0023, R2=1.00. (f) BA. Equation of line:

y=0.23x-0.0195, R2=0.99.

The value of the molar absorptivity ε was determined from the slope of the

calibration curve according to equation 3.1 and was listed in Table 3.1.

0

0.2

0.4

0.6

0.8

1

0 0.01 0.02 0.03 0.04 0.05 0.06

Concentration / mM

Ab

so

rban

ce

(e)

(f)

0

0.2

0.4

0.6

0.8

1

1.2

1.4

1.6

0 1 2 3 4 5 6 7

Concentration / mM

Ab

so

rban

ce

134

Table 3.1 Calculated molar absorptivity ε for all the modifiers studied.

Fmoc-AAs ε (L mol-1cm-1)

Fmoc-Trp 2.007 ×104

Fmoc-His 1.669 ×104

Fmoc-Phe 1.849 ×104

Fmoc-Tyr 1.901 ×104

Fmoc-Gly 1.730 ×104

BA 2.3 ×102

The surface area of CNTs was determined by measuring their specific surface

area and mass. The specific surface area of CNTs was measured by the physical

adsorption of N2 on the sample at the temperature of liquid nitrogen following

BET (Brunauer, Emmett and Teller) method 4 using a Coulter SA 3100 surface

area and pore size analyzer. During the analysis, N2 was added into an evacuated

tube containing the sample in a series of controlled doses. During this process, the

sample tube was kept at a constant temperature. The pressure in the sample vessel

was measured after the adsorption equilibrium following each dose. The recorded

pressure was used to calculate the volume of gas adsorbed which was plotted as a

function of the relative pressure (the ratio of the pressure in the sample tube to the

saturation vapor pressure of the adsorbate gas) to construct the adsorption

isotherm. The specific surface area was then calculated from the resulting

isotherm. The mass of CNTs was calculated by subtracting the mass of the Si

substrate from that of the arrays which was measured prior to introduction to the

surfactant solution. The mass of the Si substrate was determined after scratching

the CNTs off the substrate. The morphology of CVD grown CNT arrays and

randomly aligned CNT networks were characterized by SEM.

3.2.2.2 Adsorption of the surfactants on randomly aligned

CNT networks

Randomly aligned CNT networks (Figure 3.2d) were also used as a model surface

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to study the adsorption behavior of the modifiers. The substrate was prepared by

dispersing 0.2 mg of CNTs in EtOH followed by depositing onto a silicon wafer

of 1 cm x 1 cm. The dried substrate was placed into a quartz cuvette containing

3 mL of 0.05 mM Fmoc-Trp solution. UV-Vis measurements were carried out

over time to monitor the change in the concentration of the solution.

3.2.2.3 Desorption of the surfactants from CNT arrays in

H2O

The desorption behaviour of Fmoc-Trp and Fmoc-Phe from CNT arrays was

investigated in H2O. The functionalized arrays from the adsorption study were

placed into 3 mL of H2O and the UV-Vis spectra were recorded over time.

3.2.2.4 Freundlich adsorption isotherm

An adsorption isotherm is an equilibrium relationship between the quantity of the

adsorbate on the surface of an adsorbent and the equilibrium concentration of the

adsorbate in solution. There are two well established types of adsorption

isotherms: the Langmuir adsorption isotherm and the Freundlich adsorption

isotherm. The Langmuir isotherm is applicable to ideal sorption where monolayer

is formed. Whereas, the Freundlich isotherm model is applicable to non-ideal

sorption on heterogeneous surfaces where multilayers can be formed. The

Freundlich adsorption isotherm is an empirical equation based on the distribution

of the solute between the solid phase (the surface of an absorbent) and aqueous

phase at equilibrium, which is given by:

Q = k C1/n (n>1) (3.2)

Where, for the current purpose Q (mol kg -1) is the equilibrium loading of

Fmoc-AAs on unit weight of CNT arrays and C (mol m-3) is the equilibrium

concentration of Fmoc-AAs in solution. The constant k is effectively the

adsorption capacity or adsorption power of the surface for the modifiers, which

indicates the number of binding sites available on the surface. The constant n is

related to the lateral interaction between the adsorbed molecules and the

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heterogeneity of the surface which prevents this interaction, i.e. the adsorption

intensity. Equation 3.2 is linearized into logarithmic form for data fitting as

follows:

ln Q = ln k + 1/n × ln C (3.3)

Thus the constant k and n can be calculated from the intercept and slope of the

linear plot of ln Q against ln C respectively according to equation 3.3.

In the current study, to obtain the adsorption isotherm of the modifiers on CNT

arrays, a series of Fmoc-Trp solutions with concentrations of 0.025, 0.033, 0.044

and 0.051 mM were prepared. CNT arrays of known mass were placed into 3 mL

of each of solution and the UV-Vis absorption was measured over time until there

was no change. The same measurement was also conducted with Fmoc-Gly and

BA solutions.

3.2.2.5 Competitive binding from the Fmoc-AAs library on

graphite

Graphite substrates were placed into 150 mL of the library solution consisting of

5 Fmoc-AAs at a concentration of 0.08 mM (Figure 3.4). Aliquots of the solution

was regularly analysed by reversed-phase HPLC (UltiMate® 3000 Intelligent LC

system with Acclaim® 120 silica-based reversed-phase columns and UVD 170U

detector, DIONEX) to determine the concentration of each species in the solution

until no further change in UV absorbance was observed. The absorbance was

recorded at 265 nm. The HPLC chromatogram of the individual Fmoc-AA

species involved as well as that of the library solution prior to addition of the

substrate was initially recorded for reference.

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Figure 3.4 Schematic illustration of the competitive binding from the library

solution of Fmoc-AAs on graphite.

3.2.2.6 Switchable surface chemistry

In order to verify that the adsorbed Fmoc-AAs with a lower binding energy could

be displaced by those with a higher binding energy, a switchable surface was

demonstrated. Highly ordered pyrolytic graphite (HOPG) was selected as the

model surface since it is atomically flat and therefore provide a homogenous

surface for accurate contact angle measurement. A freshly cleaved HOPG surface

was obtained by peeling off the topmost layers using scotch tape.

The change in the surface chemistry was followed by water contact angle

measurement (KRÜSS drop shape analysis system DSA100). For the

measurement, a sessile drop of water (10 µL) was deposited on the sample

surface and the profile of the droplet was recorded and analysed with DSA1 v1.9

drop shape analysis software 5. The contact angle was measured on the

assumption that the droplet was at rest on the sample surface and reached

equilibrium quickly. The contact angles at both sides of the droplet were

measured and averaged. The measurements were repeated at 3 different locations

on each of the surfaces and averaged.

The contact angles on HOPG control surface and those functionalized with

Fmoc-Trp Fmoc-Phe Fmoc-Tyr Fmoc-His

Fmoc-Gly

Graphite

HPLC

UV-Vis Filtered

138

Fmoc-Trp and Fmoc-Gly were initially measured for reference. The

functionalized surfaces were prepared by placing the HOPG substrate into 3 mL

of 0.04 mM Fmoc-Trp and Fmoc-Gly solution respectively for 2 days. For the

displacement study, a freshly cleaved HOPG surface was initially placed in 3 mL

of 0.04 mM Fmoc-Gly (low binding energy) solution for 2 days followed by

drying under room temperature for contact angle measurement. The substrate was

then introduced to 3 mL of 0.04 mM Fmoc-Trp (high binding energy) solution for

2 days after which the same measurement was carried out.

3.2.3 Synthesis of CNT-inorganic nanohybrids

3.2.3.1 Synthesis of silica coated Fmoc-AA functionalized

CNTs

The catalyzing effect of the Fmoc-AAs on the sol-gel formation of silica was

initially tested without CNTs. 0.2 mL of TEOS solution (volume ratio of

TEOS:H2O:EtOH=2:1:4) was added to 15 mL of 0.04 mM Fmoc-Trp solution as

well as to H2O as a control. The mixture was sonicated for 30 min to avoid phase

separation followed by stirring for 3 days. Fmoc-His and Fmoc-Tyr solutions

were also tested under the same condition.

Subsequently, the Fmoc-AA functionalized CNTs were employed as templates

for the synthesis of CNT-silica nanohybrids via an in-situ sol-gel process at room

temperature and neutral pH (pH 7.6). In a typical experiment, 2 mg of CNTs were

dispersed in 15 mL of 0.05 mM Fmoc-Trp aqueous solution for 30 min using a tip

sonicator with a 20 s on/off pulse rate (Branson digital sonifier 250). The tip was

placed into the solution approximately one-third of the distance from the surface.

Operating in pulsed mode allows better temperature control than continuous

mode by retarding the rate of temperature increase in the medium. Additionally,

as surfactants are present, the suspension could foam during sonication, which

will interfere with the delivery of ultrasonic energy to the suspension. Pulse mode

operation allows the dissipation of the foam during the off periods. A control

dispersion in which no modifiers were added was also prepared. For the sol-gel

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coating process, a mixture of TEOS, H2O and EtOH (0.2 mL, volume ratio of

TEOS:H2O:EtOH= 2:1:4) was slowly dropped into the above CNT dispersion and

sonicated for 30 min to avoid phase separation, followed by stirring at room

temperature. The products were collected at 3 days and 21 days by centrifugation

at 8000 rpm for 1 h followed by thorough wash with H2O and ethanol for 3 times

respectively before finally dried at 90 ºC. The same procedure was performed

using Fmoc-His and Fmoc-Tyr as modifiers. The morphology of the as-produced

nanohybrids was characterized by SEM and EDX equipped with SEM was

performed to study the elemental composition of the product. TEM analysis was

conducted on samples prepared for 3 d and 21 d. The dried products were heated

in air at 200 ºC for 2 h using a ramp rate of 1 ºC/min. The annealed samples were

characterized by TEM.

The elemental distribution across a hybrid NT was measured by EDX line scan

using Tecnai F30 operating at 300 kV equipped with an EDX spectrometer. The

line scan was recorded by rastering the electron beam in a line perpendicular to

the tube axis.

The kinetics for the growth of silica coating was studied by both NMR

spectroscopy and SEM. 29Si NMR has been reported previously as a useful tool

for studying the sol-gel process of silica at the molecular level 6,7. It allows

effective identification of numerous silicate species, from dimers to prismatic

hexamers, present in aqueous silicate solutions. In the current study, the 29Si

NMR spectra were obtained on a Bruker 400-MHz NMR spectrometer at room

temperature. The acquisition time was 1.02 s and a relaxation delay of 5 s was

used. Each spectrum consisted of 128 scans. The sample for NMR measurement

was prepared by mixing 0.8 mL of TEOS with 0.4 mL of Fmoc-His/D2O solution

and 1.6 mL of EtOH (MTEOS=1.4 M) under stirring. Samples were collected at

regular intervals until 7 days and examined by NMR. The NMR spectra were

internally referenced to a tetramethylsilane (TMS) standard.

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The change in the thickness of SiO2 coating over reaction time was also

monitored by SEM to investigate the kinetics. It should be noted that this

approach does not follow silica polymerization but rather infers from the

thickness of the silica coating. The samples collected at 3, 7 and 21 days were

subjected to SEM analysis.

3.2.3.2 Synthesis of TiO2 coated Fmoc-AA functionalized

CNTs

CNT-TiO2 nanohybrids were prepared using the same surface modifiers as those

used for the silica coating experiment via the in-situ sol-gel process. Two

concentrations of CNTs (12wt% and 30wt% with respect to the expected mass of

TiO2) were used to study the role of nanotube concentration. Hybrids containing

30wt% CNTs were prepared by dispersing 2 mg of CNT in 15 mL of EtOH by

ultrasonication for 30 min, to which 1 mL of 0.05 mM Fmoc-Trp solution was

added under stirring at 0 ºC. 0.2 mL of 10% (v/v) TBOT/EtOH solution was then

added dropwise to the CNT dispersion whilst stirring. Similarly, 0.6 mL of 10%

(v/v) TBOT/EtOH solution was added to prepare the hybrids containing 12wt %

CNTs. The reaction was allowed to proceed for 1 h at 0 ºC. The products were

collected by vacuum filtering for the hybrids prepared with 12wt% CNTs, while

for those produced with 30wt% CNTs, the products were collected by

centrifugation. The products were then washed with EtOH for 3 times before

dried in air at room temperature. The same procedure was performed using

Fmoc-His, Fmoc-Tyr and BA as modifiers. A control experiment in which

pristine CNTs were used as templates was also conducted.

The morphology of the as-produced nanohybrids containing 30wt% CNTs was

characterized with SEM and EDX equipped with SEM was performed to study

the elemental composition of the products. The products prepared with both CNT

concentrations were characterized by TEM, SAED and XRD.

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To investigate the effect of CNT to modifier ratio on the morphology of the

resultant hybrids, 10s diluted modifier solution was used. Briefly, 2 mg of CNTs

were dispersed in 15 mL of EtOH, to which 1 mL of 0.005 mM Fmoc-His

solution was added, followed by mixing with 0.2 mL of 10% (v/v) TBOT/EtOH

solution for 1 h at 0 ºC. The morphology of the product was characterized by

SEM.

Similar to the study of the growth of silica, the kinetics for the growth of TiO2

coating on CNTs was studied by recording SEM images of the hybrids for

different reaction times. 2 mg of CNTs were dispersed in 15 mL of EtOH, to

which 1 mL of 0.05 mM Fmoc-Trp solution was added, followed by mixing with

0.2 mL 10% (v/v) TBOT/EtOH solution under stirring at 0 ºC. Samples were

collected after reaction for 10 min, 1 h and 6.5 h for SEM characterization.

The as-produced hybrids were heated in air at 400 ºC for 2 h to induce

crystallization of the coating followed by at 550 ºC for 2 h to oxidatively remove

CNT templates using a ramp rate of 20 ºC/min. The morphology of the calcined

samples was characterized by SEM and TEM. EDX was performed to confirm

the removal of CNT template. SAED and XRD analysis was conducted to

examine the crystal structure and phase composition of the calcined samples.

XRD measurement was performed using a Philips Automatic Powder

Diffractometer (APD), with Cu Kα radiation, 50 kV/40 mA, λ= 1.5406 Å. The

scans were performed with a step size of 0.05° (2θ) and the collection time at

each step was 25 s.

The transition from anatase to rultile was induced by heat treatment in argon at

temperatures between 800 ºC and 900 ºC. Factors including heating temperature,

pre-treatment and ramp rate were investigated for their influence on phase

transformation.

(1) To study the effect of heating temperature, the as-produced hybrids were

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heated in argon at 900 ºC and 800 ºC respectively for 2 h using a ramp rate of

20 ºC/min followed by calcination at 550 ºC for 2h.

(2) To study the effect of pre-treatment, the as-produced hybrids were initially

heated in air at 400 ºC for 2 h to induce crystallization of the coating to

anatase followed by heating in argon at 800 ºC for 2 h using a ramp rate of

20 ºC/min.

(3) To study the effect of ramp rate, the same heat treatment was performed as

that in (2) but with 1 °C/min ramp rate.

The thermally treated samples were characterized by TEM, XRD and SAED.

In a typical experiment, the aligned CNT arrays (on a 5 mm x 5 mm Si wafer)

were initially functionalized with the modifiers by immersing into 3 mL of

0.05mM Fmoc-Trp solution for 13 days to reach adsorption equilibrium. The

functionalized arrays were then placed into 0.4 mL of 10% (v/v) TBOT/EtOH

solution to which 1 mL of EtOH was added. The reaction was allowed to proceed

for 1 h at 0 ºC, after which the arrays were collected and left to dry in air. Similar

procedure was performed using Fmoc-His, Fmoc-Tyr and BA as surface

modifiers. A control experiment was also performed in which pristine CNT arrays

were coated in TBOT/EtOH solution. The dried sample was heated in air at

400°C for 2 h followed by at 550 ºC for 2 h using a ramp rate of 20 °C/min to

produce TiO2 NT arrays. The morphology of the thermally treated products was

characterized by SEM and TEM. XRD was conducted to study the crystal

structure and phase composition of the calcined samples.

3.2.3.3 Combined sites

Two dispersions were prepared by dispersing 1 mg of CNTs in 30 mL in either

0.05 mM Fmoc-His or 0.05 mM Fmoc-Tyr solution. For the single population

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experiments, a dispersion was deposited onto Si wafer and left to dry in air before

SEM characterization. For preparation of the combined catalyst, the two

dispersions were mixed with equal volume on a roller mixer for 1h before being

deposited on Si wafer. The dried sample was also observed under SEM.

For sol-gel coating of SiO2, the coated silicon wafers were immersed in a TEOS

solution (volume ratio of TEOS:H2O:EtOH=2:1:4) for 6 d followed by wash in

EtOH. The dried product was annealed in air at 200 ºC for 3 h followed by at

650ºC for 2 h to remove the carbonaceous template. For sol-gel coating of TiO2,

the combined catalyst on silicon wafers were immersed in a 10% (v/v)

TBOT/EtOH solution for 1 h at 0 ºC followed by wash in EtOH. The dried

sample was annealed in air at 400 ºC for 2 h followed by at 550 ºC for 2 h. The

morphology of the samples was characterized by SEM and EDX was performed

on both the as-produced and annealed samples.

3.2.4 Graphene and graphene based nanocomposites and

nanohybrids

3.2.4.1 GO-Inorganic nanohybrids

3.2.4.1.1 Preparation of aqueous dispersion of GO

Aqueous dispersion of GO (1 mg/mL) was produced by dispersing the

as-produced graphite oxide powder (synthesized according to a modified

Hummers method 1) in H2O using an ultrasonic bath. The bath sonicator was

fitted with a cooling system to maintain a constant temperature of 25 ºC. The

obtained suspension was centrifuged at 3000 rpm for 1 h. Only a negligible

amount of unexfoliated particles was observed after centrifugation, leading to a

homogenous dispersion of GO with a concentration of ~1 mg/mL. The obtained

GO dispersion was used for the subsequent synthesis of GO-TiO2 and GO-SiO2

nanohybrids.

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The morphology of GO was characterized by SEM and TEM. For SEM imaging,

highly diluted GO dispersion was deposited on Al stub and the dried sample was

imaged using an accelerating voltage of 10 kV. The corresponding SAED pattern

was also recorded.

3.2.4.1.2 Preparation of GO-TiO2 nanohybrids

GO-TiO2 hybrids were synthesized following an in-situ sol-gel process. In a

typical process, 0.4 mL of 1% (v/v) TBOT/EtOH solution was added dropwise to

8 mL of 1 mg/mL GO dispersion under vigorous stirring at room temperature.

The reaction was allowed to proceed for 4 h and 7 d respectively. The product

was collected by centrifugation and washed with H2O and EtOH before dried at

room temperature. Following the same procedure, hybrids with higher TBOT

concentration was also prepared by mixing 0.1 mL of 10% (v/v) TBOT/EtOH

solution with 8 mL of 1 mg/mL GO dispersion for 4 h. The morphology of the

as-produced hybrids was characterized by TEM and the corresponding SAED

patterns were recorded. EDX equipped with TEM was performed on the hybrids

prepared for 4 h with lower TBOT concentration.

The hybrids prepared with both TBOT concentrations for 4 h were heated in Ar at

500 °C for 2 h using a ramp rate of 20 ºC/min to induce TiO2 crystallization and

to avoid oxidation of GO. The annealed samples were studied by TEM and

SAED.

3.2.4.1.3 Preparation of GO-SiO2 nanohybrids

GO-SiO2 hybrids were synthesized using tetrakis (2-hydroxyethyl) orthosilicate

(THEOS) as water-soluble silica precursor via the in-situ sol-gel process. The

structure of THEOS was shown in Figure 3.5. For the sol-gel process, 15 mL of

1mg/mL GO dispersion was mixed with 150 mg of THEOS under stirring for 1

day followed by standing for 1 week. The product was collected by centrifugation

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and washed with H2O. The as-produced hybrids were characterized by SEM,

TEM, EDX and SAED. The dried product was heated in air at 600 °C for 2 h

using a ramp rate of 10 ºC/min to remove GO. The thermally treated sample was

studied by TEM and SAED.

Figure 3.5 Molecular structure of THEOS.

3.2.4.2 bwGO-Inorganic nanohybrids

3.2.4.2.1 Preparation of bwGO dispersion

For the preparation of bwGO dispersion, 15 mg of bwGO powder was dispersed

in 0.05 mM Fmoc-Trp solution at a concentration of 1 mg/mL for 10 h using an

ultrasonic bath. The obtained suspension was subjected to centrifugation at

3000rpm for 1 h. The top 80% supernatant was collected for the subsequent

synthesis of bwGO-TiO2 nanohybrids. bwGO dispersed in H2O was also prepared

as a control. Optical images of both dispersions were recorded over time for a

period of 35 d to study their stability. The morphology of bwGO was

characterized by TEM and the corresponding SAED pattern was recorded.

3.2.4.2.2 Synthesis of bwGO-TiO2 nanohybrids

(1) Reaction in aqueous solution

0.4 mL of 1% (v/v) TBOT/EtOH solution was mixed with 11 mL of bwGO

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dispersion whilst stirring for 4 h. The product was collected by centrifugation and

washed with EtOH.

(2) Reaction in EtOH

1 mg of bwGO was initially dispersed in 10 mL of EtOH followed by adding

1mL of 0.05 mM Fmoc-Trp solution. The mixture was further sonicated for

10min before mixing with 1 mL of 1% (v/v) TBOT/EtOH solution for 4 h. The

product was collected by centrifugation and washed with EtOH. A control

experiment in which 1 mL of H2O was added instead of Fmoc-Trp solution was

also conducted.

The products were characterized by SEM, EDX, TEM and SAED. The dried

products from both reactions were annealed in Ar at 600 ºC for 2 h to induce

crystallization of TiO2 for Raman characterization. Raman spectra for bwGO and

anatase TiO2 were also recorded for reference.

3.2.4.3 Exfoliated graphene (EG)-Inorganic nanohybrids

3.2.4.3.1 Preparation of graphene dispersion

A set of graphene dispersions were prepared by direct exfoliation of graphite in

Fmoc-Trp solution using a range of sonication times and centrifuge speeds. The

procedure is similar to that reported for using sodium dodecyl benzene sulfonate 8

and sodium cholate as surfactants 9. In a typical procedure, graphite powder

(grade 2369) was added to 30 mL of 0.05 mM Fmoc-Trp solution to give an

initial concentration of 3 mg/mL. The mixture was sonicated in a bath sonicator

for 1 h. As the sonic energy input to the sample is sensitive to the exact position

of the sample in the bath and the water level 10, all the samples were sonicated in

a fixed position in the bath. To minimize water evaporation as a result of the

temperature increase for prolonged sonication, the bath sonicator is fitted to a

cooling system to stabilize the temperature of the bath at 25 ºC. A black liquid

consisting of a homogenous phase and large numbers of shiny graphite particles

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was obtained after sonication. The suspension was subjected to centrifugation at

3000 rpm for 1 h. The top 80% supernatant was decanted by pipette for the

subsequent analysis. The same procedure was performed for the preparation of

graphene dispersions under other conditions as listed in Table 3.2. Digital images

of the dispersions were recorded and their stability was studied by monitoring the

change in homogeneity of the dispersion over a period of 1 week.

Table 3.2 Conditions used for the preparation of graphene dispersions

Dispersion Initial concentration

(mg/mL)

Sonication time

(h)

Centrifuge speed

(rpm)

1 3 1 3000

2 3 6 3000

3 3 12 3000

4 3 6 500

5 3 6 6000

In order to determine the exact concentration of the dispersion, a relatively large

amount of freshly prepared dispersion of known volume was vacuum filtered

through an alumina membrane of known mass. The film formed on the membrane

was washed with H2O to remove residual modifiers followed by drying in

vacuum oven overnight at room temperature. The mass of the dried film was then

determined using a microbalance.

3.2.4.3.2 Preparation of EG-TiO2 nanocomposites and

nanohybrids

(1) Preparation of EG-TiO2 nanocomposites in aqueous solution

0.2 mL of 1% (v/v) TBOT/EtOH solution was mixed with 11 mL of graphene

dispersion prepared with 6 h of sonication followed by centrifugation at 6000 rpm

and 3000 rpm respectively under sonication for 30 min followed by stirring for 4h.

The product was collected by centrifugation and washed with EtOH.

(2) Preparation of EG-TiO2 nanohybrids in EtOH

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1.1 mg of graphite powder was initially dispersed in 15 mL of EtOH for 2 h,

followed by adding 1 mL of 0.05 mM Fmoc-Trp solution to stabilize the

dispersion. The mixture was further sonicated for 6 h before mixing with 1 mL of

1% (v/v) TBOT/EtOH solution. After a further sonication for 2 h, the reaction

was allowed to proceed for another 2 h under stirring. The product was collected

by centrifugation and washed with EtOH.

The products obtained from both reactions were characterized by TEM, SAED

and Raman spectroscopy.

3.2.5 Mineralization of peptide self-assembled hydrogels

3.2.5.1 Fmoc-Y hydrogel preparation

Fmoc-Tyrosine (phosphate)-OH (Fmoc-Y(p)-OH) was dissolved in 3 mL of 0.6M

phosphate buffer at a concentration of 40 mM and the pH was adjusted to neutral.

To this solution, alkaline phosphatase (10 DEA μL-1) was added and the

self-supporting hydrogel of Fmoc-Tyr was formed after the solution had been

kept at 37 ºC for 2 h.

3.2.5.2 Fmoc-FY hydrogel preparation

For preparing 2 mL of gel, 12.6 mg of Fmoc-FpY was dissolved into 1 mL of

phosphate buffer to which 10 μL of 1 M HCl was added. The mixture was

vortexed for 30 s before made up to 2 mL by adding buffer. 200 μL of AP solution

of 3 DEA was added to the Fmoc-FpY solution followed by vortex for 20 s. The

resulting solution was left for gelation at room temperature after 90 min.

3.2.5.3 Characterization

The structure of the hydrogels was characterized by TEM. The sample was

negatively stained prior to TEM observation to improve image contrast. The

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formvar-carbon coated grids were initially glow discharged followed by placing

the grid formvar side down on a drop of the hydrogel (~1-3 μL) for 10 s. Excess

gel was removed by placing the grid on a drop of distilled water (~10 μL) for 60 s

followed by blotting with filter paper. The washed grid was stained by placing on

a 10 μL droplet of 2% (w/v) uranyl acetat for 60 s followed by blotting with filter

paper.

The conversion of Fmoc-FpY into Fmoc-FY by enzymatic dephosphorylation

was followed using reverse phase HPLC. Aliquots of the sample solution were

collected at time points: t=0, 5, 10, 30 and 90 min and was mixed with an

acetronitrile/water mixture (50:50) containing 0.1% trifluoroacetic acid for HPLC

analysis.

Fluorescence spectroscopy was carried out to examine the supramolecular

arrangement within the self-assembled Fmoc-FY fibres. Fluorescence emission

spectra of both the Fmoc-FY hydrogel and the solution of Fmoc-FpY were

measured on a Jasco FP-6500 spectrofluorometer with excitation at 295 nm.

3.2.5.4 Silicification of Fmoc-Y gel

Two different methods were performed for the silicification of Fmoc-Y gel.

Method 1: 4 μL of TEOS was vortexed in the diluted Fmoc-Y hydrogel (8 mM)

for 30 s to allow homogeneous distribution of the precursor. The sol-gel process

was then allowed to proceed for different periods under ambient conditions prior

to TEM analysis. The composition of the mineralized sample was studied by

EDX.

Method 2: A mixture of TEOS in water was added onto the top of a volume of

Fmoc-Y hydrogel at 40 mM (volume ratio of hydrogel:TEOS:H2O=1:1:1) and

was allowed to stand at room temperature. After aging for 1 month, two phases

separated by a white layer were observed. The white layer was characterized by

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TEM and EDX. Both the upper aqueous phase and the lower hydrogel phase were

characterized by SEM.

3.3 Analytical techniques

3.3.1 Scanning Electron Microscopy (SEM)

SEM is a microscopy technique that images the sample surface by scanning it

with a focused beam of high-energy electrons in a raster scan pattern. A variety of

signals are produced from the interaction of the electrons with the atoms at or

near the surface of the sample, which contain information about the sample's

surface topography and composition. Among the generated signals, secondary

electron imaging shows the higher resolution of the surface topographical

characteristics than backscattered electron imaging. SEM can produce very

high-resolution images of a sample surface, revealing details less than 1 nm in

size.

In the current study, a field emitter gun scanning electron microscope (FEGSEM,

Philips XL30) was employed to study the morphology of the samples. Unless

otherwise specified, an accelerating voltage of 20 kV was used. Samples for SEM

analysis were prepared by depositing sample dispersion in EtOH onto the

polished aluminium stub (Agar Scientific) and leave to dry in air. In order to

image the CNT arrays, the sample was mounted onto the holder using carbon tape

with the axis of the tube perpendicular to the incident electron beam. For

nonconductive samples, a thin coating of Pt or C was sputtered on the sample

using a Precision Etching Coating System (PECS, Gatan) to reduce charging

effect. After coating, silver paint was applied to form a conductive “bridge” from

the sample surface to the sample holder. The obtained micrographs were analyzed

using ImageJ software.

3.3.2 Transmission Electron Microscopy (TEM)

http://en.wikipedia.org/wiki/Microscope

151

In a TEM instrument, a beam of electrons is emitted from an electron gun and is

focused to a thin and coherent beam by the condenser lenses. The beam is

restricted by the condenser aperture, knocking out high-angle electrons, before

interacting with the sample as it passes through. The transmitted beam is focused

and enlarged by the objective and projector lens respectively and form an image

on the phosphor screen. Figure 3.6 illustrates a schematic diagram of a TEM.

TEM offers two modes of operation: image mode and diffraction mode. The

diffraction mode is used to study the structure and composition of crystals. The

difference between the two modes may only be the strength of the intermediate

lens (Figure 3.7). Instead of focusing on the first image plane in image mode,

the diffracted beams are brought to a focus in the back focal plane of the

objective lens in diffraction mode. In imaging mode, an objective aperture is

inserted in the back focal plane to enhance contrast by blocking out high-angle

diffracted electrons (Figure 3.7a). For selected area electron diffraction (SAED), a

selected area aperture in the plane of the first intermediate image defines the

region of which the diffraction is obtained (Figure 3.7b). The contrast observed in

the TEM image may be brought about by several mechanisms including

mass-thickness contrast, diffraction contrast and phase contrast.

http://ncmn.unl.edu/cfem/about/glossary.shtml#aperture

http://ncmn.unl.edu/cfem/microscopy/interact.shtml

http://ncmn.unl.edu/cfem/microscopy/interact.shtml

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Figure 3.6 Schematic diagram of a TEM. 11

Figure 3.7 Ray path in a TEM operating in (a) image mode (b) diffraction mode. 12

In the present study, a Philips CM200 TEM operating at 200 kV was used to

characterize the morphology of the samples. For TEM observation, the samples

were prepared by depositing a droplet (~10 μL) of the sample dispersion in EtOH

onto a holey-carbon coated Cu grid (Agar scientific, 400 mesh) followed by

solvent evaporation in air. The obtained micrographs were analyzed using ImageJ

software. The corresponding SAED patterns were also acquired.

To determine the dimensions of the exfoliated flakes, the sample dispersion was

deposited onto holey carbon coated Cu grid for TEM analysis. The flake size

distribution was studied by measuring 40 flakes from each of the dispersions and

(b) (a)

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the flake area was calculated using AxioVision Rel. 4.8 software. SAED analysis

was performed on the sample prepared with 6h/500 rpm.

3.3.3 Energy Dispersive X-ray Spectroscopy (EDX)

EDX is a microanalytical technique that uses the characteristic spectrum of

X-rays emitted by the specimen after excitation by high-energy electrons to

obtain information about its elemental composition. EDX is typically used to

detect elements of high atomic number. The spatial resolution is determined by

the probe size, beam broadening within the specimen, and the effect of

backscattered electrons on the specimen around the point of analysis.

In the current study, the elemental composition of the samples was analyzed with

an energy-dispersive X-ray spectrometer equipped with SEM (FEGSEM, Philips

XL30) operating at 10 kV.

3.3.4 Reversed-phase high-performance liquid

chromatography (RP-HPLC)

An HPLC system consists of a stationary phase (column) and a mobile phase

(solvent). When the sample is injected into the system, it is carried by the mobile

phase flowing through the stationary phase. The different affinities between the

sample components to the mobile phase and to the stationary phase result in the

separation. The sample component attracting to the stationary phase has a longer

retention time, thus coming out of the column at the last while those attracting to

the mobile phase come out of the column very quickly. When a non-polar

material is chosen to be the stationary phase, the HPLC system is called

RP-HPLC. When using this technique, the component with the highest

hydrophobicity has the longest retention time (Figure 3.8).

154

Figure 3.8 Schematic representation of reversed-phase HPLC. The most hydrophilic

components (orange) elute from the column first, followed by the less hydrophilic

components (green), and finally the most hydrophobic components (blue). 13

3.3.5 Contact angle measurement

The contact angle θ is defined as the angle formed between the liquid-solid

interface and the liquid-vapor interface which is acquired by applying a tangent

line from the contact point along the liquid-vapor interface in the droplet profile

(Figure 3.9). Unlike ideal surfaces, a drop placed on a non-ideal surface has a

spectrum of contact angles ranging from the advancing contact angle ( ) to the

receding contact angle ( ). The equilibrium contact angle ( ) is within those

values, and can be calculated from and . Contact angle hysteresis is

defined as the difference between the advancing and receding contact angles.

155

Figure 3.9 Schematic of a liquid drop on a solid surface, where the

solid–vapor interfacial energy is denoted by γsv, the solid–liquid interfacial energy

is denoted by γsl, and the liquid–vapor interfacial energy is denoted by γlv. 14

The contact angle quantifies the wettability of a solid surface by a liquid. A

contact angle less than 90° indicates that wetting of the surface is favorable, such

as hydrophilic surfaces; while contact angles greater than 90° generally means

that wetting of the surface is unfavorable such as hydrophobic surfaces.

The contact angle is determined using the sessile drop method (Figure 3.10). A

droplet of water is deposited on the sample surface through a syringe needle. The

profile of the drop is recorded with a camera and analysed using DSA1 v1.9 drop

shape analysis software 5. The baseline which is defined as the projection of the

sample surface in the drop image is first determined automatically. A contour

recognition is carried out based on a grey-scale analysis of the image. A

geometrical model describing the drop shape is then fitted to the contour. The

contact angle is given by the angle between the calculated drop shape function

and the sample surface using circle fitting method.

156

Figure 3.10 Sessile drop method for determining the contact angle. The fitted

contour is shown in green. 15

3.3.6 Raman spectroscopy

3.3.6.1 Background

Raman spectroscopy has been historically used to characterize the structure and

quality of graphitic materials 16,17 and it has recently been proved as a reliable and

nondestructive tool for capturing the electronic structures of graphene 18.

Raman spectrum for defect-free graphitic materials typically shows two

characteristic features (Figure 3.11a): G band at ~1580 cm-1 and 2D band at

~2700 cm-1 with 514 nm laser excitation. The G band originates from a single

resonance process associated with the doubly degenerate (iTO and LO) phonon

modes (E2g symmetry) at the Brillouin zone center 19,20. The 2D band which is

historically named G’ band is the second order of D band. The 2D band is

associated with two phonon inter-valley double resonance scattering involving

iTO phonon near the K point20. In defected graphitic materials, a defect-induced

D band is observed at ~1350 cm-1 with 514 nm laser excitation (Figure 3.11b).

157

The D band originates from a second-order double resonance Raman process

involving inter-valley scattering of iTO phonon near the K point of the Brillouin

zone 21. The defects are generally divided into edge defects 22 and structural

defects on the basal plane. The D peak intensity is proportional to the amount of

disorder 21. The intensity ratio of D to G band, ID/IG, is generally used to quantify

the defect content. Double resonance can also happen as intra-valley process, i.e.

connecting two points belonging to the same cone around the K or K’ points,

which gives rise to the so-called D' peak at ~1620 cm-1

in defected graphite 23.

Figure 3.11 (a) Typical Raman spectra for bulk graphite and monolayer graphene

obtained using a 514 nm laser. (b) Comparison of the D band at 514 nm at the edge

of bulk graphite and monolayer graphene. The fit of D1 and D2 components of the D

band of bulk graphite is shown. 18

Ferrari et al. 18 have suggested that the number of layers in few-layered graphene

(<5 layers) with AB stacking can be identified by the position, lineshape and

linewidth of 2D band. For monolayer graphene, the 2D band is fitted by a single

Lorentzian peak with a FWHM of ~30 cm-1 24, whereas it splits into four and six

components in bi- and trilayer graphene respectively 20 (Figure 3.12). Few-layer

graphene with more than 4 layers showed a more complex and broader 2D peak,

which is difficult to be distinguished from that of HOPG. For graphite, the 2D

band is described with two components, 2D1 and 2D2, which are roughly 1/4 and

1/2 the height of the G peak respectively. The distinct feature of monolayer is that

the 2D band in monolayer graphene is approximately 2 to 4 times more intense

G 2D

158

than that of the G band 20.

Figure 3.12 Measured 2D band for (a) monolayer, (b) bilayer, (c) trilayer, (d)

four-layer and (e) HOPG using a 514 nm laser. 20

3.3.6.2 Raman characterization of the exfoliated samples

The degree of exfoliation and the quality of the exfoliated flakes was investigated

using Raman spectroscopy. Raman spectra for the exfoliated samples and the

starting graphite powder were recorded on a Renishaw 2000 Raman spectroscopy

system (Renishaw Instruments, England) using a 633 nm HeNe excitation laser.

The laser beam was focused on the sample using a ×50 objective lens. A low

laser power of ~1mW was used to avoid laser induced sample heating. Spectra

were collected with an exposure time of 20 s and accumulation of 3 times.

Samples for Raman characterization were prepared by depositing the dispersion

onto SiO2/Si substrate (with 300 nm thick SiO2 layer) which was heated at 110 °C

on a hotplate to accelerate solvent evaporation. The SiO2/Si substrates were

159

cleaned by sonication in acetone followed by in H2O and isopropanol. At least 40

flakes deposited from each of the samples were randomly picked and measured

for statistical study. The spectra were fitted using a Lorentzian function to

determine the parameters of the bands using WiREInterface software.

3.3.7 Atomic Force Microscopy (AFM)

AFM is a type of scanning probe microscopy (SPM), providing a 3D profile of

the surface with a resolution on the order of fractions of a nanometer.

A typical AFM system consists of a cantilever with a sharp tip at its end (with a

typical radius of a few to tens of nm) which is used to scan the sample surface.

When the tip is brought into proximity of a surface, forces between the tip and the

sample lead to a deflection of the cantilever according to Hooke's law. The

deflection is typically measured using a laser beam reflected off the top surface of

the cantilever into a position-sensitive photodiode detector (Figure 3.13). The

measured deflections are used to generate a map of the surface topography.

There are 3 primary imaging modes in AFM: contact mode, non-contact mode

and tapping mode (intermittent-contact mode or Dynamic Force Mode (DFM)).

In contact mode, the cantilever tip remains in contact with the sample surface

during scanning and the contours of the surface are obtained by maintaining a

constant cantilever deflection employing a feedback mechanism. In non-contact

mode, the cantilever is oscillated at or near its natural resonance frequency. The

detection scheme is based on monitoring changes in the resonant frequency or

amplitude of the cantilever using a feedback loop. The non-contact mode is

preferable to contact mode for measuring soft samples. The tapping mode

combines qualities of both modes by oscillating the cantilever tip at or near its

natural resonance frequency while allowing the tip to touch or tap the sample

surface intermittently during scanning. The advantage of tapping the surface is

the improved lateral resolution on soft samples.

http://en.wikipedia.org/wiki/Scanning_probe_microscopy

http://en.wikipedia.org/wiki/Nanometer

http://en.wikipedia.org/wiki/Cantilever

http://en.wikipedia.org/wiki/Force

http://en.wikipedia.org/wiki/Hooke%27s_law

http://en.wikipedia.org/wiki/Laser

http://en.wikipedia.org/wiki/Scanning_probe_microscopy#Constant_interaction_mode

160

Figure 3.13 Schematic diagram of the beam deflection system in an atomic force

microscope, using laser and photodetector to measure the beam position. 25

AFM measurement was performed on the exfoliated samples to study the degree

of exfoliation. The measurement was taken with a VEECO CPII Atomic Force

Microscope working in contact mode.

Samples for AFM measurements were prepared by depositing a drop (~100 μL)

of highly diluted graphene dispersion onto a freshly cleaved mica surface (Agar

Scientific) which was heated at ~110 °C on a hotplate to accelerate solvent

evaporation. AFM image of Fmoc-Trp solution was also recorded as a control.

All the images were analyzed using IP 2.1 software.

In order to accurately identify the number of layers per flake, it is necessary to

remove the modifiers from the flakes. Two approaches were attempted: (1) to

rinse the sample in H2O for 5 min followed by blowing dry under a stream of

161

argon (2) to anneal the sample in Ar at 300 ºC for 1 h. AFM images of the treated

samples were recorded.

3.4 Reference

1. W. S. Hummers Jr. et al., Preparation of Graphitic Oxide, J. Am. Chem. Soc.,

1958, 80, 1339.

2. J. P. Rourke et al., The Real Graphene Oxide Revealed: Stripping the Oxidative

Debris from the Graphene-like Sheets, Angew. Chem. Int. Ed., 2011, 50, 3173.

3. C. Singh et al., Production of aligned carbon nanotubes by the CVD injection

method, Physica B, 2002, 323, 339.

4. S. Brunauer et al., Adsorption of Gases in Multimolecular Layers,

J. Am. Chem. Soc., 1938, 60, 309.

5. DSA v1.9 Drop Shape Analysis Manual, p65-68.

6. L.W. Kelts et al., Sol-gel chemistry studied by 1H and 29Si nuclear magnetic

resonance, Journal of Non-Crystalline Solids, 1986, 83, 353.

7. J. C. Pouxviel et al., NMR study of the sol/gel polymerization, Journal of

Non-Crystalline Solids, 1987, 89, 345.





6, 458.


2010, 6, 864.

11. http://ncmn.unl.edu/cfem/microscopy/TEM.shtml

12. http://www.microscopy.ethz.ch/TEMED.htm

13. http://www.atdbio.com/content/7/Purification-of-oligonucleotides

14. http://www.ramehart.com/contactangle.htm

15. http://www.kruss.de/services/education-theory/glossary/drop-shape-analysis/ 16. M. A. Pimenta et al., Studying disorder in graphite-based systems by Raman

spectroscopy, Phys.Chem.Chem.Phys., 2007, 9, 1276.

17. M. S. Dresselhaus et al., Raman spectroscopy of carbon nanotubes, Phys.

Rep., 2005, 409, 47.

18. A. C. Ferrari et al., Raman Spectrum of Graphene and Graphene Layers, Phys.

Rev. Lett., 2006, 97, 187401.

19. F. Tuinstra et al., Raman Spectrum of Graphite, J. Chem. Phys., 1970, 53,

1126.

20. L.M. Malard et al., Raman spectroscopy in graphene, Phys. Rep., 2009, 473,

51.

21. A. C. Ferrari et al., Interpretation of Raman spectra of disordered and

amorphous carbon, Physical Review B, 2000, 61, 14095.

22. C. Casiraghi et al., Raman Spectroscopy of Graphene Edges, Nano Lett., 2009,

http://www.kruss.de/services/education-theory/glossary/drop-shape-analysis/

162

9, 1433.

23. R. J. Nemanich et al., First- and second-order Raman scattering from

finite-size crystals of graphite, Physical Review B, 1979, 20, 392.

24. Y. F. Hao et al., Probing Layer Number and Stacking Order of Few-Layer

Graphene by Raman Spectroscopy, Small, 2010, 6, 195.

25. http://en.wikipedia.org/wiki/Atomic_force_microscopy.

http://en.wikipedia.org/wiki/Atomic_force_microscopy

163

Chapter 4 Dynamic Interaction of Fmoc-AAs with

CNTs

4.1 Introduction The majority of the research on the noncovalent functionalisation of CNTs using

biomolecules has tended to focus on the response of CNTs. For example, the

concentration of CNTs dispersed in solution is measured but the binding kinetics

of the biomolecules on the nanotubes is not considered. However, the interaction

energy is equally important as if it is low then the surfactants will be desorbed

from the CNTs’ surfaces when placed in excess solvent, causing the dispersion to

crash. Furthemore, it is crucial to investigate the adsorption behaviour of the

biosurfactants on CNTs’ surfaces before employing them as templates for

producing hybrid materials. To address this question, the dynamic adsorption of a

series of aromatic Fmoc-AAs on CNTs was investigated 1. The binding of the

Fmoc-AAs was considered initially by measuring the interactions of single

species with CNTs. The Freundlich isotherm model was used to describe the

adsorption isotherm. However, these experiments took considerable laboratory

time to identify the Fmoc-AA with the highest affinity for the nanotubes.

Therefore, a competitive binding approach was developed to screen efficiently a

library of Fmoc-AAs for the strongest binding candidate. Finally, a switchable

surface chemistry was demonstrated to verify that Fmoc-AAs with a higher

binding energy could displace those with a lower binding energy on the surface.

4.2 Synthesis of aligned MWNT arrays by injection

CVD method Aligned MWNTs arrays were successfully grown on oxidised silicon substrates

through an injection CVD method using ferrocene as catalyst and toluene as

carbon source. The ferrocene-toluene solution was injected into an argon carrier

gas preheated to 200 ºC, where they were vaporized and passed into the heating

164

zone of the furnace (760 ºC). Previous study has found that the synthesis

temperature of 760 ºC produced the nanotubes with the least defects 2. Fe catalyst

particles were produced by the decomposition of ferrocene and deposited on the

SiO2 layer of the substrate to catalyze the growth of CNTs.

The surface morphology of the synthesized CNTs was characterized using SEM.

The aligned CNT arrays showed an average length of ~430 μm (Figure 4.1a). A

magnified SEM image revealed generally good alignment of “wavy” CNTs along

their length, with the presence of some entanglements (Figure 4.1b). TEM

analysis of the CNTs found their diameter in the range of 22-97 nm, with an

average value of 56.6 nm and a standard deviation of 16.6 nm (Figure 4.1c). The

dark particles (with an average diameter of a few nanometers) entrapped either

within the hollow cavity or the walls of the NTs (indicated by arrows in Figure

4.1c) represented the Fe catalyst residues. The HRTEM image revealed the

multi-walled nature of the CNT which showed a relatively clean surface (Figure

4.1d). Only a negligible layer of amorphous carbon was observed to deposit on

the graphitic layers. The corresponding SAED pattern (Figure 4.1e) exhibited a

bright arc for the (002) diffraction together with a ring for the (100) diffraction,

and a faint arc for the (004) diffraction of MWNTs.

165

Figure 4.1 SEM images of CNT arrays grown at 760 ºC from a 5wt% ferrocene in

toluene solution on SiO2 substrate for 1h. (a) Cross-sectional image of the aligned

CNT arrays. (b) Close-up view of the CNTs from the arrays. (c) TEM image of the

pristine CNTs with dark particles presented both in the hollow cavity and the walls

of CNTs (indicated by arrows). Scale bar, 0.2 μm. (d) HRTEM image showing the

multilayered structure of a synthesized CNT with the lattice fringes clearly visible.

Scale bar, 5 nm. (e) The corresponding SAED pattern was indexed to the (002), (100)

and (004) planes of MWNTs.

200 μm

430 μm

(a)

5μm

(b)

(c

)

(d) (002)

(e) (004)

(100)

(002)

166

4.3 Interaction of surface modifiers with CNTs

An effective and well-studied interface between the biosurfactants and CNTs is

essential to their application in biology as well as the performance of CNT based

hybrid materials. To this end, the adsorption behaviour of a library of Fmoc-AAs

on CNTs as a function of the amino acid was evaluated. The four aromatic amino

acids (Trp, Phe, His and Tyr) were studied with Gly as a control as these amino

acids had shown promise as surfactants for dispersing CNTs 3. The adsorption

isotherm of benzyl alcohol (BA) was also studied as a reference surfactant as it

has been previously used in the literature as a surface modifier for the production

of CNT-TiO2 nanohybrids4. The effect of surface chemistry of CNTs on the

morphology of the synthesized hybrid materials was investigated in later

Chapters.

4.3.1 Adsorption behavior of the modifiers on aligned

CNT arrays

The synthesized CNT arrays were used as the model surface as they did not

disperse into solution hence allow the concentration of the modifiers in solution

to be measured. The adsorption process of single Fmoc-AA species on CNT

arrays was monitored by recording the UV-Vis spectra of the solution in the

presence of the arrays over time. The spectra were also collected in the absence of

the surface as a reference. Take Fmoc-Trp for example, as shown in Figure 4.2a,

the concentration of Fmoc-Trp control solution remained constant over a period

of 3 days suggesting that its precipitation and settling did not occur. In contrast, in

the presence of the CNT arrays, a significant drop in the concentration of the

Fmoc-Trp solution was observed within the first 100 min, after which the rate of

change slowed and eventually a nearly constant concentration value was reached

indicating adsorption equilibrium. Similar profiles were obtained for the other

Fmoc-AA solutions studied (Figure 4.2c,e and i) except for Fmoc-His which still

had not reached equilibrium after 6 days (Figure 4.2g). The affnity of the studied

Fmoc-AAs for CNTs was evaluated in terms of their equilibrium loading and

167

initial adsorption rate which were normalised to surface area of CNTs. Fmoc-Trp

was found to have the best affinity for CNTs by showing an equilibrium loading

of 0.000325 mmol/m2 CNT (Figure 4.2k) while the non-aromatic Fmoc-Gly had

the lowest loading of 0.000072 mmol/m2 CNT due to the lack of interaction

between the non-aromatic side chain and CNTs. The adsorption profiles confirm

that the Fmoc group was an efficient anchor, which could be improved by the

addition of an aromatic amino acid. Quantum mechanical modelling has

highlighted the importance of π-π stacking interactions in these systems 3. Having

studied the effect of the amino acid side chains on the affinity for CNTs, Fmoc

was compared with other aromatic ligands with different degree of aromaticity,

such as the benzyl group of BA. The adsorption profile of BA on CNT arrays

showed a significant increase in the concentration of BA solution after 1 day

(Figure 4.2l). This was caused by the slow evaporation of the solvent although the

UV-Vis cuvette was sealed. This result suggests the poorer affinity of BA for

CNTs which is possibly attributed to the lower degree of aromaticity of the

benzyl ring compared with that of the fluorenyl ring of the Fmoc group.

The initial adsorption rate, ki, is an important characteristic in defining the

adsorption efficiency. ki is given by the initial change in the moles of the

adsorbate on unit surface area of the adsorbent per unit time:

Sdt

VdC

Sdt

dnki (4.1)

Where, for the current purpose, n (nmol) is the amount of Fmoc-AAs adsorbed on

the arrays, S (m2) is the surface area of CNTs, C (nmol/L) is the concentration of

Fmoc-AAs in solution and V (L) is the volume of the solution. dC/dt is obtained

by drawing the tangent to the adsorption profile at the initial linear phase. Take

Fmoc-Trp for example, the adsorption rate was determined over the initial linear

region of the profile from 0 to 60 min with a coefficient of determination R2 of

0.98 (Figure 4.2b). The initial adsorption rates for the other Fmoc-AAs studied

were determined using the same method (Figure 4.2d,f,h and j) and summarized

168

in Table 4.1. All the adsorption curves were found to display linearity over the

same time range with Fmoc-Trp being adsorbed the most quickly while Fmoc-His

showing the lowest rate.

0 1 2 3 4 50.03

0.04

0.05

0.06

Fmoc-Phe control

In the presence of CNT mats

Co

nce

ntr

atio

n o

f F

mo

c-P

he

/ m

M

Time / day

(c)

0 1 2 3 4 50.03

0.04

0.05

0.06

Fmoc-Tyr control


Co

nce

ntr

atio

n o

f F

mo

c-T

yr

/ m

M

Time / day

(e)

0 1 2 3 40.03

0.04

0.05

0.06

Fmoc-Trp control

in the presence of CNT mats

Co

nce

ntr

atio

n o

f F

mo

c-T

rp /

mM

Time / day

(a)

0 20 40 60 800.044

0.046

0.048

0.050

0.052

0.054

Co

nce

ntr

atio

n o

f F

mo

c-T

rp /

mM

Time / min

(b)

0 20 40 60

0.047

0.048

0.049

0.050

0.051

Co

nce

ntr

atio

n o

f F

mo

c-P

he /

mM

Time / min

(d)

0 20 40 60 80

0.047

0.048

0.049

0.050

0.051

Concentr

ation o

f F

moc-T

yr

/ m

M

Time / min

(f)

169

Figure 4.2 Adsorption profiles of (a) Fmoc-Trp (c) Fmoc-Phe (e) Fmoc-Tyr (g)

Fmoc-His (i) Fmoc-Gly and (l) BA on aligned CNT arrays. (b,d,f,h,j) Determination

of the initial adsorption rate of the corresponding modifiers on the arrays. (k)

Histogram showing the equilibrium loadings of the Fmoc-AAs on the arrays.

0 1 2 3 40.03

0.04

0.05

0.06

Fmoc-Gly control


Co

nce

ntr

atio

n o

f F

mo

c-G

ly /

mM

Time / day

(i)

(g)

0 1 2 3 4 5 6 7

0.03

0.04

0.05

0.06

Fmoc-His control

In the presence of CNT matsC

oncentr

ation o

f F

moc-H

is /

mM

Time / day

(k)

AA = Trp Phe Tyr His Gly

0.0000

0.0001

0.0002

0.0003

0.0004

The

equili

bri

um

loadin

g o

f F

moc-A

As /

mm

ol /

m2

CN

T

Fmoc-AA

0 1 2 3 4 5 6

4.3

4.4

4.5

4.6

4.7

4.8

4.9

5.0

5.1

in the presence of CNT mats

BA control

Co

nce

ntr

atio

n o

f B

A /

mM

Time / day

(l)

0 20 40 60 80

0.0480

0.0485

0.0490

0.0495

0.0500

Co

nce

ntr

atio

n o

f F

mo

c-G

ly /

mM

Time / min

(j)

0 20 40 60 80

0.0500

0.0504

0.0508

0.0512

Concentr

ation o

f F

moc-H

is /

mM

Time / min

Equation

Weight

Residual Sum of Squares

Pearson's r

Adj. R-Square

B

B

(h)

170

Table 4.1 Initial adsorption rate of the Fmoc-AAs on CNT arrays

Initial adsorption rate

(nmol m-2 min-1)

R2

Fmoc-Trp 0.00263±0.00013 0.98

Fmoc-Phe 0.00116±0.00009 0.96

Fmoc-Tyr 0.00102±0.00003 0.99

Fmoc-His 0.00037±0.00002 0.98

Fmoc-Gly 0.00083±0.00005 0.96

4.3.2 Adsorption behavior of the modifiers on randomly

oriented CNT networks

To examine whether the modifiers could diffuse into the lower part of the arrays,

the adsorption of the Fmoc-AAs on randomly aligned CNT network was studied.

As shown in the adsorption profile (Figure 4.3), the concentration of Fmoc-Trp

solution decreased linearly over time and after 9 days of adsorption the

equilibrium was not reached. The spectra recorded after this period observed the

blue shift of the peak from 264 nm to 258 nm, possibly due to the precipitation of

the amino acid. The loading of Fmoc-Trp on the network after 9 days was

calculated to be 0.16 mmol/g CNT, while the loading on the arrays was

determined to be 0.058 mmol/g CNT, indicating that the full surface area of CNTs

in the arrays was not accessible to the modifiers.

Figure 4.3 Adsorption profile of Fmoc-Trp on randomly aligned CNT networks.

0 2 4 6 8 10

0.038

0.040

0.042

0.044

0.046

0.048

0.050

0.052

Co

nce

ntr

atio

n o

f F

mo

c-T

rp / m

M

Time / day

171

4.3.3 Desorption behavior of the modifiers in excess of

water

To establish if the interaction between Fmoc-AAs and CNTs are dynamic, the

desorption behaviour of Fmoc-Trp and Fmoc-Phe from CNT arrays was studied

in excess of H2O. As shown in Figure 4.4, the desorption profiles of both

Fmoc-AAs revealed a new adsorption equilibrium after ～50 h with the desorbed

Fmoc-AAs accounted for ~27% of the overall Fmoc-AAs previously adsorbed on

the arrays.

Figure 4.4 Desorption profiles of (a) Fmoc-Trp and (b) Fmoc-Phe from CNT arrays

in water.

This result verifies the reversible nature of the binding process of Fmoc-AAs

which were in dynamic equilibrium between solution and adsorption on CNT

arrays (Equation 4.2). The adsorption equilibrium is a function of temperature, as

at higher temperatures more Fmoc-AAs will exceed their binding energy and

enter solution.

T Fmoc-AAs in solution

T

↓

↑ (Desorption dominate) (Adsorption dominate) (4.2) Fmoc-AAs adsorbed on CNT arrays

0 20 40 60 80 100 120

0

10

20

30

40

Pe

rce

nta

ge

of F

mo

c-T

rp d

eso

rbe

d / %

Time / h

Model

Equation

Reduced Chi-Sqr

Adj. R-Square

B

B

B

B

B

B

B

Model

Equation

Reduced Chi-Sqr

Adj. R-Square

B

B

B

(a)

0 50 100 150 200

0

10

20

30

40

Pe

rce

nta

ge

of

Fm

oc-P

he d

eso

rbe

d /

%

Time / h

(b)

172

4.3.4 Freundlich isotherm model

The previous TEM analysis has shown that Fmoc-AAs absorbed heterogeneously

on CNTs and tended to form spherical aggregates, ranging from 5 to 10 nm in

diameter 3. Therefore, adsorption isotherm such as the Langmuir isothermal

which describes the formation of a monolayer is inappropriate for the Fmoc-AA

system. Instead, the Freundlich isotherm model which is applicable to non-ideal

sorption on heterogeneous surfaces where multilayers can be formed was used.

The adsorption isotherms of Fmoc-Trp, Fmoc-Gly and BA on CNT arrays were

studied. As shown in Figure 4.5, the equilibrium data for both Fmoc-Trp and

Fmoc-Gly were found to follow the Freundlich isotherm model over the initial

concentration range between 0.025-0.055 mM with a coefficient of determination,

R2, of the linear regression greater than 0.90. While all the adsorption experiments

of BA showed an increased concentration of the solution over time, similar to that

observed in section 4.3.1. This result further confirmed that Fmoc was a more

effective ligand than the benzyl ring of BA due to the higher degree of

aromaticity of the fluorenyl ring.

The values of the constants k and n for Fmo-Trp and Fmoc-Gly were determined

from the plot and given in Table 4.2. It was found that the value of k for

Fmoc-Trp was 25 times that for Fmoc-Gly, indicating the higher adsorptive

capacity of CNTs for Fmoc-Trp. The molecular modeling study has shown the

greater interaction energy of ionized Fmoc-Trp with SWNTs (-47.7 kcal/mol)

over ionized Fmoc-Gly (-36.2 kcal/mol)3. Furthermore, it was found that the

aromatic side chain in Trp encouraged the flattening of Fmoc-Trp against the NT

by following its curvature, which favors the formation of multilayer films,

whereas the side chain of Gly lifted away from the NT surface. Both the higher

binding energy and flat conformations of Fmoc-Trp on CNTs’ surface

contributed to the higher adsorption capacity.

For Freundlich isotherm, n>1 represents favourable adsorption conditions and

173

thus the calculated values of n for both Fmoc-AAs suggested that they were

favourably adsorbed on CNT arrays. Fmoc-Gly had a greater n than Fmoc-Trp

suggesting that the Fmoc-Gly molecules interacted more between themselves

within the adsorbed layer.

Figure 4.5 Plot of ln Q vs. ln C for the adsorption of Fmoc-Trp (red circles) and

Fmoc-Gly (blue triangles) on the arrays.

Table 4.2 Calculated adsorption capacity (k) and intensity (n) for Fmoc-AAs

adsorbed on CNT arrays. Note that the units for k depend on the value of n. The

quality of fit, R2, was also given for each Fmoc-AA.

Fmoc-AAs k n R2

Fmoc-Trp 5.65 m3 kg-1 1.0 0.90

Fmoc-Gly 0.23 mol1/3 m2 kg-1 1.5 0.98

4.3.5 Competitive binding from the Fmoc-AAs library on

graphite

Although the affinity of the Fmoc-AAs for CNTs could be evaluated by

measuring the adsorption profile of each species, however this method is very

time consuming. Therefore, if a solution containing a library of Fmoc-AAs was

used instead of a singular species, it is expected that there would be competition

-4.5 -4.0 -3.5 -3.0 -2.5 -2.0

-5

-4

-3

-2

-1

Fmoc-Trp

Fmoc-Gly

ln Q

ln C

174

between the species to absorb onto the NTs’ surface.

The competitive binding from the Fmoc-AAs library on graphitic surfaces was

assessed using graphite as model surface by reversed-phase HPLC. The retention

time for the individual Fmoc-AA species was measured for reference (Figure

4.6a-e). The HPLC chromatogram of the library solution was then recorded prior

to introducing the graphite substrates (Figure 4.6f). The HPLC traces of the

mixture were recorded over time until the adsorption equilibrium was reached. It

was found that out of all the library components, only the concentration of

Fmoc-Trp became significantly lower when equilibrium was reached (173 h),

whereas the other Fmoc-AAs showed little change in their concentrations (Figure

4.7a). These results suggested that Fmoc-Trp had come out of solution and

predominately bound to the surface of graphite, while the others with poorer

affinity would be predominantly in solution. It is proposed that the stronger

binding Fmoc-AAs would displace the Fmoc-AAs which bound more weakly

from graphite. Over time, the strongest binding Fmoc-AA would be

predominantly on the surface of graphite whereas the weaker binding Fmoc-AAs

would be predominantly in solution.

The equilibrium loading of each component from the competitive binding

experiments was calculated and compared to those from the individual adsorption

experiments. In contrast to the individual adsorption experiment in which all the

aromatic Fmoc-AAs adsorbed at a similar level 1, the competitive binding

approach clearly identified the strongest binder as Fmoc-Trp which showed 3.5

times the loading of the next strongest Fmoc-AA (Fmoc-Phe) (Figure 4.7b). Such

a competitive binding approach allows a number of binding candidates with

similar binding energies to be evaluated simultaneously for identifying the

strongest binder in a very efficient manner. The data for the individual adsorption

experiments are taken from the Master’s work 1.

175

Figure 4.6 HPLC chromatogram of 0.4 mM of (a) Fmoc-Phe (b) Fmoc-Trp (c)

Fmoc-Tyr (d) Fmoc-Gly and (e) Fmoc-His. (f) The mixture of the 5 Fmoc-AAs with

the same volume ratio.

0

200

400

600

800

14 19 24 29 34

Time / min

UV absorbance / a.u.

His Gly

Tyr

173 h

0 h

Trp Phe

(a)

0

2000

4000

6000

8000

10000

14 19 24 29 34

Time / min

UV

ab

so

rba

nce

/ a

.u.

Mixture of 5 Fmoc-AAs

Fmoc-His-OH

Fmoc-Gly-OH

Fmoc-Tyr-OH

Fmoc-Trp-OH

Fmoc-Phe-OH

(f)

(e)

(d)

(c)

(b)

(a)

176

Figure 4.7 (a) HPLC traces of the mixture consisting of the five Fmoc-AAs at 0 h

(upper) and after 173 h of competitive binding (lower). (b) Comparison of the

equilibrium loadings of the five Fmoc-AAs on graphite in individual adsorption and

competitive binding experiments.

4.3.6 Switchable surface chemistry

In order to verify that the Fmoc-AAs with a higher binding energy could displace

those with a lower binding energy on the surface, a switchable surface chemistry

was demonstrated through sequential exposure of HOPG surface to Fmoc-Gly

(low binding energy) and Fmoc-Trp (high binding energy). The change in the

surface chemistry was followed using contact angle measurement. The contact

angles on control HOPG surface and those immobilized with Fmoc-Trp and

Fmoc-Gly were initially measured for reference. As shown in Figure 4.8a, The

contact angles on control HOPG surface was 89.8 ° ± 2.8 ° indicating its

hydrophobic nature whereas those on Fmoc-Trp (Figure 4.8b) and Fmoc-Gly

(Figure 4.8c) functionalised HOPG surfaces were 76.1° ± 0.7 ° and 53.3 ° ± 1.6 °

respectively. The decrease in the contact angle suggested that after

Fmoc-Trp-OH Fmoc-Phe-OH Fmoc-Tyr-OH Fmoc-His-OH Fmoc-Gly-OH

0.00

0.04

0.08

0.12

0.16

Th

e e

qu

ilib

riu

m lo

ad

ing

of F

mo

c-A

As

(in

div

idu

al a

dso

rptio

n)

/ m

mo

l/m2 g

rap

hite

equilibrium loading in individual adsorption

euqilibrium loading in competitive binding

0

1

2

3

4

5

Th

e e

qu

ilib

riu

m lo

ad

ing

of F

mo

c-A

As

(co

mp

etit

ive

bin

din

g)

/ m

mo

l/m2 g

rap

hite

(b)

177

immobilization with Fmoc-Trp and Fmoc-Gly, the surface become more

hydrophilic due to the presence of the carboxylic groups on the surface. The

lower contact angle for Fmoc-Gly functionalized HOPG surface was due to that

Fmoc-Gly was more hydrophilic than Fmoc-Trp.

The clear difference in contact angles between the two functionalized surfaces

was then utilized to examine the displacement of Fmoc-Gly by Fmoc-Trp on the

surface. After introducing Fmoc-Gly functionalized HOPG surface to Fmoc-Trp

solution for 2 days, the contact angle increased from 53.3° ± 1.6 ° to 77.1° ± 1.1 °,

which was similar to that on the Fmoc-Trp functionalised HOPG surface,

suggesting that Fmoc-Gly was partially displaced by Fmoc-Trp from the HOPG

surface.

178

θ = 89.8 ° ± 2.8 ° θ = 53.3 ° ± 1.6 ° θ = 77.1 ° ± 1.1 °

Figure 4.8 Displacement of Fmoc-Gly by Fmoc-Trp on HOPG surface. Contact

angle measurements on (a) freshly cleaved HOPG surface (b) Fmoc-Trp

functionalized HOPG surface and (c) Fmoc-Gly functionalized HOPG surface. (d)

Schematic representation of the displacement of Fmoc-Gly by Fmoc-Trp on the

surface. (e-g) Corresponding contact angle measurements on the surfaces presented

in (d). Note that Figure (c) and (f) were repeats of the same experimental condition.

4.4 Conclusion

The adsorption behavior of a series of aromatic Fmoc-AAs on aligned CNT

HOPG

(a)

θ= 89.8 º ± 2.8 º

HOPG + Fmoc-Trp

(b)

θ= 76.1 º ± 0.7 º

(c)

HOPG + Fmoc-Gly

θ= 53.3 º ± 1.6 º

(d)

Fmoc-Trp Fmoc-Gly

(e)

HOPG HOPG + Fmoc-Gly

(f)

After displacement

(g)

179

arrays was studied. The adsorption kinetics and equilibrium of single Fmoc-AA

species on the surface were measured using UV-Vis spectroscopy. Among the

Fmoc-AAs studied, Fmoc-Trp was found to have the best affinity for CNTs by

showing the highest equilibrium loading and initial adsorption rate, whilst the

non-aromatic control Fmoc-Gly showed the lowest equilibrium loading. These

results confirmed that the Fmoc group was an efficient anchor, which could be

improved by the addition of an aromatic amino acid. The adsorption behaviour of

the Fmoc-AAs on randomly aligned CNT networks was also studied and the

result indicated that the modifiers could not diffuse into the lower part of the

arrays, thus could not access the full surface area of the nanotubes. The fully

reversible nature of the binding process was demonstrated via the desorption of

the modifier from CNTs’ surface in excess of water.

The equilibrium data were well fitted to the Freundlich isotherm model over the

initial concentration range between 0.025-0.055 mM. Both the higher binding

energy and flat conformations of Fmoc-Trp on CNTs’ surface contributed to the

higher adsorption capacity compared with Fmoc-Gly.

The competitive binding from the library of Fmoc-AAs on graphite was

developed. Fmoc-Trp was identified as the strongest binding candidate at the

expense of the other components studied, leading to a significantly different

binding behavior compared with the individual adsorption experiments. It is

proposed that the stronger binding Fmoc-AAs would displace the Fmoc-AAs

which bound more weakly from graphite. This approach provides an efficient

way to screen a wide range of binding candidates with similar binding energies

simultaneously for the strongest binder.

A switchable surface chemistry was demonstrated through the sequential

displacement of Fmoc-Gly by Fmoc-Trp from a HOPG surface. This observation

supported the hypothesis that the Fmoc-AAs with a higher binding energy could

displace those with a lower binding energy from the surface.

180

This study on the dynamic interaction between these aromatic amino acid

derivatives and CNTs not only provides a step forward for their bioapplication

where an effective and well-studied interface is required but also pave the way

towards the subsequent utilization of these functionalized CNTs as templates for

the production of hybrid materials.

4.5 References

1. Y. Li et al., A study of the dynamic interaction of surfactants with graphite and

carbon nanotubes using Fmoc-amino acids as a model system, Langmuir, 2009,

25, 11760.

2. C. Singh et al., Production of controlled architectures of aligned carbon

nanotubes by an injection chemical vapour deposition method, Carbon, 2003, 41,

359.



4. D. Eder et al., Carbon–inorganic hybrid materials: the carbon-nanotube/TiO2

interface, Adv. Mater., 2008, 20, 1787.

181

Chapter 5 Synthesis of CNT-inorganic

nanohybrids and the corresponding inorganic NTs

using Fmoc-AAs as surface modifier

5.1 Introduction

Inorganic coatings are typically coated onto CNTs using sol-gel processes due to

the benign reaction conditions used (eg. room temperature, near neutral pH etc.).

The sol-gel process also allows control over the morphology and properties of the

resultant inorganic networks through adjusting the reaction parameters, such as

the reaction time1 and choice of metal precursor 2. Eder et al. 3 have employed

BA as a surfactant to coat pristine CNTs with TiO2 via a sol-gel process. They

assumed that the benzene ring of the surfactant adsorbed on CNTs’ surface via

- stacking interactions, whilst the hydroxyl groups contributed to the

hydrolysis of TBOT and further induced condensation to form a Ti–O–Ti

network.

The Fmoc-AAs investigated in Chapter 4 can be considered as surfactants similar

to BA, but with the Fmoc group providing better binding to the CNTs than the

one-membered benzyl ring, and the AAs providing a range of potential surface

chemistries. Having previously established the interaction of the Fmoc-AAs with

CNTs, three were selected as surface modifiers for sol-gel coating: Fmoc-Trp,

Fmoc-His and Fmoc-Tyr. These Fmoc-AA functionalized CNTs were then used as

templates for the production of silica and titania based nanohybrids via the sol-gel

process. The Fmoc-AA choice was influenced by previous reports that both

imidazole4,5 and hydroxyl groups2,3,6 could catalyze SiO2 and TiO2 deposition

from solution. The morphology of the nanostructures was characterized using

SEM and TEM. The effect of AA on the morphology of the deposited coating

was also investigated. The kinetics for the growth of the inorganic layers was

studied by SEM. Anatase TiO2 NTs were obtained after calcination of the

CNT-TiO2 nanohybrids to remove the carbonaceous template. The major

advantage of using templates is that the dimension of the synthesized NTs can be

182

easily controlled by the size of the templates. The influence of heating

temperature, pre-treatment and ramp rate on the crystallization and phase

transformation was also investigated by TEM, XRD and ED.

Aligned TiO2 NT arrays have been previously synthesized by anodization of

titanium thin films 7,8 and atomic layer deposition (ALD) of TiO2 within a porous

alumina membrane 9. However, such methods either require specialized setup or

suffer from the limited length of the obtained arrays. To the best of my

knowledge there has been no report on the sol-gel templating of TiO2 NT arrays

using aligned CNT arrays as templates. The challenge is to coat individual CNTs

while preserving the morphology of the CNT structures. In the present study, a

simple route towards the synthesis of aligned TiO2 NT arrays was demonstrated

using CVD grown CNT arrays as templates in the presence of the same surface

modifiers.

5.2 Synthesis of CNT-silica nanohybrids using

Fmoc-AAs as surface modifier

5.2.1 Synthesis and morphology characterization

The possibility of the Fmoc-AAs acting as catalysts for the sol-gel formation of

silica was initially investigated by adding TEOS solution to Fmoc-Trp, Fmoc-His

and Fmoc-Tyr solution respectively as well as to H2O as a control. In the

presence of the Fmoc-AAs, precipitation was observed after reaction for 3 days,

indicating the formation of silica structures. In absence of the Fmoc-AAs, the

reaction mixture was still clear after the same period of time. This result suggests

that the Fmoc-AAs are catalysts for the sol-gel synthesis of silica.

The Fmoc-AA functionalized CNTs were then employed as templates for the

synthesis of CNT-silica nanohybrid materials at room temperature and neutral pH

(pH 7.6) via the in-situ sol-gel process. Tetraethyl orthosilicate (TEOS) was used

as the silica precursor. In absence of the Fmoc-AAs, TEOS were found to

183

randomly nucleate to form separate silica particles without attaching to the CNTs

which showed a smooth surface with an average diameter of 56.0 ± 10.1 nm

(Figure 5.1a). This lack of interaction was due to there being no affinity between

the pristine CNTs and silica precursors. In contrast, in presence of Fmoc-Trp and

Fmoc-His, a relatively uniform layer of coating was observed on the individual

CNTs as evidenced by the presence of a rougher surface (Figure 5.1b and c). Free

SiO2 particles were rarely present in the samples. As for Fmoc-Tyr functionalized

CNTs, a mixture of partially coated and uncoated CNTs were observed (Figure

5.1d). This observation suggests the importance of the surfactants in controlling

the morphology of the final hybrid materials. The elemental composition of the

product was analyzed using energy-dispersive spectroscopy. The presence of

elemental signals of C,O and Si confirmed the existence of SiO2 coating on CNTs

(Figure 5.1e). The Al signal was originated from the SEM stub. The quantitative

analysis revealed that the atomic ratio of O:Si was ~4. The rather high value may

be due to the presence of a very thin layer of Al2O3 on the top of the aluminum

substrate or the EDX analysis software assuming depth, take-off angles and

absorption values for a solid planar surface.

500nm

(b)

500nm

(a)

silica particle uncoated

500nm

184

Figure 5.1 SEM images of (a) the product obtained from the control experiment in

which pristine CNTs were used as templates. (b) Silica coated Fmoc-Trp and (c)

Fmoc-His functionalized CNTs. (d) A mixture of partially coated and uncoated

CNTs in the presence of Fmoc-Tyr after reaction for 21 days. (e) EDX spectrum of

the product shown in (c). Note that the aluminum and some of the oxygen were

from the sample stub.

In order to study the change in the silica coating morphology over reaction time,

TEM images were taken on samples which had been coated for periods between 3

to 21 days. In absence of Fmoc-AAs, silica precipitated in solution in form of

large clusters and the nanotubes remained uncoated (Figure 5.2a). In contrast, a

relatively uniform layer of silica was found on Fmoc-Trp (Figure 5.2b and c) and

Fmoc-His functionalized CNTs (Figure 5.2d and e) respectively. There were

several difficulties in accurately measuring the thickness of SiO2 coating.

Although the diameter of SiO2 coated CNTs can be determined based on SEM

Element at. %

C 57.77

Al 35.1

Si 1.12

Cu 1.57

O 4.44

(e)

0 2 4 6

0

2000

4000

6000

8000

10000

Co

un

ts (

a.u

.)

Energy (keV)

C

Cu

Al

Si O

500nm

(c)

uncoated

coated

(d)

500nm

185

images, the CNT template was not visible, making it impossible to directly

measure the thickness of the coatings. It was possible to measure the thickness in

TEM, however in bright field imaging it was difficult to determine exactly where

the interface between the SiO2 coating and the CNT was, leading to errors in the

measurement. The average coating thickness was determined to be 3.6 ± 0.6 nm

and 4.7 ± 0.5 nm for Fmoc-Trp (Figure 5.2b) and Fmoc-His (Figure 5.2d)

functionalized CNTs respectively after 3 days of reaction. Whereas, Fmoc-Tyr

functionalized CNTs were only partially covered (Figure 5.2f). No agglomeration

of the CNTs was observed during the coating process indicating that the aromatic

Fmoc-AAs were highly effective in dispersing CNTs in aqueous solution and

preventing the formation of bundles. After 21 days reaction, the silica coating on

Fmoc-Trp and Fmoc-His functionalized CNTs became thicker and less uniform

compared with those formed at 3 days, with the thickness increasing to

7.3±0.6nm (Figure 5.2c) and 6.8 ± 2.6 nm (Figure 5.2e) respectively. In contrast,

Fmoc-Tyr functionalized CNTs still exhibited a partial coating (Figure 5.2g).

These observations are in good agreement with the SEM analysis. The measured

coating thickness is summarized in Table 5.1.

186

3 days 21 days

Control

Fmoc-Trp

Fmoc-His

Fmoc-Tyr

(a)

(b) (c)

(d) (e)

uncoated

(f) (g)

uncoated

187

Figure 5.2 TEM images of (a) pristine CNTs co-existed with isolated SiO2 particles.

Note. The image was over-focused as it was taken during early stage of the PhD.

Silica coated Fmoc-Trp functionalized CNTs after reaction for (b) 3 days and (c) 21

days. Silica coated Fmoc-His functionalized CNTs after reaction for (d) 3 days and

(e) 21 days. Partially coated Fmoc-Tyr functionalized CNTs after reaction for (f) 3

days and (g) 21 days. Scale bar, (a) 100 nm, (b) 20 nm, (c)-(g) 50 nm.

Table 5.1 Measured SiO2 coating thickness based on TEM images.

Day 3 Day 21

Fmoc-Trp 3.6 ± 0.6 nm 7.3 ± 0.6 nm

Fmoc-His 4.7 ± 0.5 nm 6.8 ± 2.6 nm

Fmoc-Tyr partial coating partial coating

The homogeneity of the SiO2 coating on the CNTs was investigated by measuring

the elemental distribution across a hybrid NT using EDX line scan. Figure 5.3a

showed the line profile measured perpendicular to the tube axis direction as

indicated by the white arrow in the STEM image (inset). C, Si and O elements

were detected and the intensity ratio of the Si and O elements was constant,

indicating the presence of SiOx. It was noted that the Si and O signals showed a

maxima on the left side of the hybrid tube. This may due to the increased

interaction length of the electron beam with the edge of the coating (red line in

Figure 5.3b) compared with the center of the tube (yellow line in Figure 5.3b),

thus more x-ray photons were emitted. The absence of the maxima on the other

side of the tube was possibly caused by the drift in electron beam with time or the

thinner coating present on that side of the tube.

Reaction time

Fmoc-AAs Coating thickness

188

Figure 5.3 (a) Line profile taken perpendicular to the tube axis direction. Inset:

Dark field STEM image of the hybrid NT. The direction of the scan was marked by

the arrow. The analysis was conducted with the help of Xiaofeng Zhao. (b) Cross

sectional view of a SiO2 coated CNT. The interaction of electron beam with the edge

200 nm

0 20 40 60 80 100 120 140 160 180

0

50

100

150

200

Carbon

Silicon

OxygenC

ou

nts

(a

.u.)

position (nm)

(a)

Electron beam

(b)

189

and the centre of the hybrid tube was indicated by the red and yellow line

respectively. Blue colour: silica coating.

5.2.2 Discussion on the role of Fmoc-AA functionalization

in controlling the morphology of the hybrids

Table 5.2 Correlation of the adsorption equilibrium of the Fmoc-AAs on CNT mats

with the morphology of the hybrids

Fmoc-AAs Equilibrium loading

(mmol /m2 CNT)

Structures produced

None - No coating, only isolated SiO2 particles

Fmoc-Trp 0.000325 Uniform coating with thickness of 7.3 ± 0.6 nm

Fmoc-His 0.000301 Uniform coating with thickness of 6.8 ± 2.6 nm

Fmoc-Tyr 0.000258 Mixture of partially coated and uncoated CNTs

Table 5.2 shows the correlation of the adsorption equilibrium of the Fmoc-AAs

on CNTs with the structure of the obtained hybrids. It was found that with more

Fmoc-AAs adsorbed on CNTs, more uniform coating was observed. This could

be attributed to the presence of a lower density of binding sites on Fmoc-Tyr

functionalized CNTs for silica deposition compared with Fmoc-Trp and

Fmoc-His functionalized CNTs. It is probable that the weaker H-bonding

interaction between Tyr and hydrolyzed TEOS (O−H… :O = 21 kJ/mol) compared

with those between Trp or His and hydrolyzed TEOS (O−H… :N = 29 kJ/mol)

also contributed to the partial coating of CNTs.

5.2.3 Growing mechanism of silica coating on Fmoc-AA

functionalized CNTs

The surface modifier is proposed to play a dual role in the coating process. Firstly,

the Fmoc-AA acts as an electrostatic surfactant as under the neutral pH of the

reaction medium, the carboxylic acid group exposed on the amide backbone is

deprotonated to produce overall negative charge which stabilize the dispersion of

CNTs 10. This dispersion favours individually coated CNTs. Secondly, the

functionalities such as indole, imidazole and hydroxyl groups provides binding

190

sites for silica deposition. Previous studies 11,12 have demonstrated that the

hydroxyl or imidazole group alone was not sufficient to catalyze the hydrolysis of

TEOS. Therefore the functionalities studied in the present work were only

capable of catalyzing silicic acid condensation 4,5. The schematic illustration for

the interaction between the hydrolyzed TEOS and the Fmoc-AAs is shown in

Figure 5.4. For example, under the neutral pH of the reaction medium, the

imidazole side chain of Fmoc-His could strongly attract the hydrolyzed precursor

towards CNTs’ surface through both electrostatic (Figure 5.4a) and H-bonding

interactions (Figure 5.4b). This attraction results in an increase in the local

concentration of silica precursors in the vicinity of CNTs and consequently SiO2

nuclei are preferentially formed on the surface of CNTs through

polycondensation between adjacent precursors. The nuclei will further adsorb and

condense with the additional precursors from solution to build up of the SiO2

coating. Whereas, the indole and hydroxyl side chains of Fmoc-Trp (Figure 5.4c)

and Fmoc-Tyr (Figure 5.4d) respectively can only form H-bonds with the

hydrolyzed TEOS due to their unionizable natures at neutral pH. This observation

was in contrary to the proposed mechanism that the presence of both imidazole

side chain of histidine-165 and hydroxy group of serine-26 was required for the

efficient catalysis of silica synthesis at neutral pH 13.

(a)

(b)

191

Figure 5.4 Proposed catalytic mechanisms for silica templating. (a) Electrostatic

attraction between the protonated imidazole group of Fmoc-His and silicate anion.

H-bonding between (b) imidazole group of Fmoc-His, (c) indole group of Fmoc-Trp,

(d) hydroxyl group of Fmoc-Tyr and silanol group of hydrolyzed TEOS.

5.2.4 Kinetics for silica growth

Previous studies have reported that 29Si NMR can be used to study the sol-gel

process of silica at the molecular level 14 as it allows effective identification of

numerous silicate species, from dimers to prismatic hexamers, present in aqueous

silicate solutions. Therefore, the kinetics of silica polymerization in the presence

of Fmoc-AAs was followed with 29Si NMR. A mixture of TEOS and

Fmoc-His/D2O solution and EtOH with the volume ratio of 2:1:4 was examined

without CNTs due to the presence of magnetic Fe catalyst particles in CNTs

which would interfere with the measurement. The NMR spectrum of the 29Si

nucleus measured for the sample showed only the monomer peak after standing

for 7 days despite the same experiment conducted outside the NMR observed the

precipitation of silica after reaction for the same period of time.

(c)

(d)

192

Given the issues with the NMR approach, the change in the thickness of SiO2

coating was monitored over time by SEM to investigate the kinetics. This

approach of course does not follow the polymerization reaction but rather infers

from the thickness of the silica coating. The reaction was allowed to proceed for 3,

7 and 21 days. Due to the issue with SEM observation, the corresponding SiO2

coating thickness can not be measured. Instead, the change in the diameter of the

hybrid NT was followed by SEM and plotted against the growth time (Figure

5.5d). It was found that the average diameter of the hybrid NTs increased from

69.1 ± 14.3 nm after 3 days (Figure 5.5a) to 88.6 ± 13.5 nm after 7 days (Figure

5.5b) and to 111.4 ± 23.2 nm after 21 days of reaction (Figure 5.5c). This

observation reveals the potential of sol-gel method in controlling the thickness of

inorganic coating via tuning the reaction time.

(a) (b)

(c)

0 5 10 15 20 25

0

50

100

150

Dia

me

ter

of th

e h

yb

rid

NT

/ n

m

Growth time / day

(d)

193

Figure 5.5 SEM images of silica coated Fmoc-His functionalized CNTs obtained

after a growth time of (a) 3 days (b) 7 days and (c) 21 days. (d) Plot of the diameter

of the hybrid NT against the growth time. The average value was calculated based

on 50 separate measurements.

5.2.5 Annealing

The hybrid NTs were annealed in air at 200 °C for 2 h. TEM images of the treated

samples revealed more uniform and compact coverage of silica on CNTs (Figure

5.6b and d) as compared with those observed for the as-produced hybrids (Figure

5.6a and c), possibly due to the further crosslinking between the non-bridged

silanol groups in the coating at high temperature.

Before annealing

After annealing

Fmoc-Trp

Fmoc-His

Figure 5.6 TEM images of silica coated Fmoc-Trp functionalized CNTs (a) before

and (b) after annealing at 200 °C, and silica coated Fmoc-His functionalized CNTs

(c) before and (d) after annealing under the same condition. Scale bar, 50 nm.

(a) (b)

(c) (d)

194

5.3 Synthesis of CNT-TiO2 nanohybrids using

Fmoc-AAs as surface modifier

5.3.1 Synthesis and morphology characterization

The same Fmoc-AAs that were used for the silica coating experiments described

above were also used as surface modifiers for the synthesis of CNT-TiO2

nanohybrids. Titanium butoxide (TBOT) was used as titania precursor and all the

reactions proceed under 0 °C for 1h. The concentration of CNTs was 30 wt% with

respect to the expected mass of TiO2. The morphology of the products was

studied by SEM. A control experiment was run in which pristine CNTs were used

as templates. Similar to the synthesis of CNT-SiO2 nanohybrids, large amounts of

isolated TiO2 particles were observed to randomly co-exist with uncoated CNTs

in the resulting mixture as evidenced by the diameter of the NTs which closely

matched that of the pristine CNTs (Figure 5.7a). In contrast, uniform coating was

observed on individual CNTs in the presence of Fmoc-Trp (Figure 5.7b),

Fmoc-His (Figure 5.7c) and Fmoc-Tyr (Figure 5.7d) as evidenced by the increase

in the diameter of the NTs to 78.3 ± 17.4 nm, 86.3 ± 14.3 nm and 104.5 ±16.0 nm

respectively. BA was also used as the surface modifier for comparison. A partial

and thin coating was deposited on CNTs in the presence of BA, along with the

presence of isolated TiO2 particles (Figure 5.7e). This observation suggests the

weaker π–π stacking interaction between BA and CNTs compared with that of the

aromatic Fmoc-AAs, possibly attributed to the lower degree of aromaticity of BA,

which was consistent with the adsorption study in Chapter 4. It is worth noting

that the uniform TiO2 coating on BA functionalized CNTs observed in Eder’s

study 2,3 was not achieved in the current study due to the addition of considerably

lower amount of BA to make the results comparable. Based on these results, it is

concluded that the surface chemistry of CNTs played a key role in controlling the

morphology of the hybrids. EDX analysis confirmed the presence of C, Ti and O

in all of the synthesized nanohybrids (Figure 5.7f-h). Quantitative EDX analysis

revealed that the atomic ratio of O:Ti was ~2.4, which was slightly higher than

that of TiO2, possibly due to the same issue as discussed in Section 5.2.1.

195

0 2 4 60

500

1000

1500

2000

2500

3000

3500

Co

un

ts (

a.u

.)

Energy (keV)

C

O

Al

Ti

(h)

Cu

Element at. %

Ti 2.86

Cu 1.65

Al 32.6

C 57.07

O 5.82

(e)

1μm

1μm

(d)

1μm

(c)

(b)

1μm

(a)

0 2 4 60

500

1000

1500

2000

2500

3000

3500

Co

un

ts (

a.u

.)

Energy (keV)

C O

Al

Pt Ti

(g)

Cu

0 2 4 60

500

1000

1500

2000

2500

3000

3500

Co

un

ts (

a.u

.)

Energy (keV)

C

O Pt Ti

Al

(f)

Cu

196

Figure 5.7 SEM images of (a) the product obtained using pristine CNTs as

templates. TiO2 coated CNTs in the presence of (b) Fmoc-Trp (c) Fmoc-His (d)

Fmoc-Tyr and (e) BA. (f-h) EDX spectra measured for the hybrids shown in (b-d).

Note the Al signal was originated from SEM stub, and Pt signal was originated from

the conductive coating on the SEM sample to reduce charging effect. The

considerably stronger C signal in (h) was due to the application of a thin layer of

carbon on the SEM sample as the conductive coating.

5.3.2 Mechanism for the formation of TiO2 coating on

the functionalized CNTs

Fmoc-AAs played a similar role in the sol-gel coating of TiO2 on CNTs as

proposed for the synthesis of CNT- SiO2 nanohybrids. A possible mechanism for

the formation of TiO2 coating on Fmoc-AAs functionalized CNTs was discussed

as follows: In absence of Fmoc-AAs, due to the hydrophobic nature of the

pristine CNTs’ surface there was no attraction with the titanium precursor,

therefore the precursors prefer to nucleate and precipitated as nanoparticles in

solution. In the presence of the Fmoc-AAs, the functionalities (such as indole,

imidazole and hydroxyl groups) exposed on the surface of CNTs attracted the

hydrolyzed TBOT or TiO2 nuclei through H-bonding interactions, and thus

promoted the condensation between adjacent precursors to form Ti–O–Ti

network on CNTs. Whether hydrolyzed TBOT or TiO2 nuclei are deposited on

the CNTs would depend on the rate of reaction upon introduction of TBOT into

the system and the ratio of TBOT to water. In order to follow the reaction rate,

DLS would be performed in further work. Park et al. 15 have reported that for a

complete hydrolysis of TBOT, a H2O:Ti ratio of 4.8 was needed. In the current

study, for the hybrids produced with 12 wt% of CNTs, the ratio of H2O:Ti was

~32 and for those produced with 30 wt% of CNTs, the ratio of H2O:Ti was ~95.

Therefore, for all the systems studied TBOT should have been fully hydrolyzed.

The proposed formation process of CNT-TiO2 hybrids and the subsequent TiO2

NTs was illustrated in scheme 5.1. In the presence of BA, although the benzene

ring facilitated the π-π stacking on CNTs’ surface, the weaker interaction

compared with that between the fluorenyl ring of Fmoc-AAs and CNTs led to

197

lower number density of binding sites on CNTs for TiO2 deposition. This

observation indicated the important role of the aromatic Fmoc-AAs in the

formation of the hybrid nanostructures. Presumably, TiO2 coating is not formed

via a Ti–O–C bond, as is the case for acid-treated CNTs, but rather via π-π

stacking interaction between CNTs and the fluorenyl ring of Fmoc-AAs 3.

Scheme 5.1 Schematic illustration of the preparation of CNT-TiO2 nanohybrids and

TiO2 NTs in the presence of Fmoc-His. The exposed imidazole groups on CNT’s

Calcination

TiO2 NT TiO2 coated CNT

TBOT

Hydrolysis

Condensation

Fmoc-His

198

surface promoted the interaction with hydrolyzed titanium precursors through

H-bonding interaction.

5.3.3 Effect of CNT to TBOT ratio on the hybrid

morphology

In order to investigate the effect of the concentration of CNTs on the morphology

of the produced hybrids, two different concentrations of CNTs, i.e. 12 wt% and

30 wt% were used. The structure of the products was characterized by TEM. As

shown in Figure 5.8a, in the control experiment where pristine CNTs were used

as templates, TiO2 was found to form nanoparticle aggregates without coating the

CNTs. By contrast, relatively uniform TiO2 coatings on individual CNTs were

observed in the presence of Fmoc-Trp, Fmoc-His and Fmoc-Tyr for both CNT

concentrations (Figure 5.8b-g). It was noted that in the presence of 12 wt% of

CNTs, the TiO2 coating on CNTs was occasionally broken. This may be induced

by the drying stress involved during the vacuum filtration process which led to a

fast removal of solvents and thus to considerable contraction of the coating. For

samples produced with the CNT concentration of 30 wt%, the products were

separated from the reaction media by centrifugation, therefore, no cracks were

observed on the coating. In the case of using BA as the surface modifier, a

negligible layer of coating was observed in the presence of 30 wt% of CNTs

(Figure 5.8h) while partially coated CNTs were resulted in the presence of 12wt%

of CNTs (Figure 5.8i). The observation is in good agreement with the SEM study.

In addition, the coating thickness was found to be highly dependent on the ratio

of CNT to TBOT. The thickness of the TiO2 coating was calculated based on the

TEM observation and summarized in Table 5.3. These results demonstrated that

the sol-gel coating thickness can be readily controlled by adjusting the ratio of

CNT to the precursor. The corresponding SAED pattern of the as-produced

hybrid (Figure 5.8j) exhibited a bright arc for the (002) diffraction together with a

ring for the (100) diffraction, and a faint arc for the (004) diffraction of MWNTs.

The absence of any diffraction typical for either anatase or rutile suggested that

the TiO2 coating at this stage was amorphous. XRD analysis further confirmed

199

the amorphous nature of the as-produced TiO2 coating. Figure 5.8k showed the

(002) diffractions of C (2θ=26.2°) superimposed by a broad background which

indicated the presence of an amorphous phase.

30 wt% CNTs 12 wt% CNTs

Control

Fmoc-Trp

Fmoc-His

(b) (c)

(d) (e)

(a)

200

Figure 5.8 TEM images of (a) the product obtained using pristine CNTs as

templates. TiO2 coated CNTs in the presence of Fmoc-Trp with the CNT

concentration of (b) 30 wt% and (c) 12 wt%. TiO2 coated CNTs in the presence of

Fmoc-His with the CNT concentration of (d) 30 wt% and (e) 12 wt%. A cluster of

TiO2 nanoparticles were deposited on the smooth surface of the TiO2 coating in (e).

TiO2 coated CNTs in the presence of Fmoc-Tyr with the CNT concentration of (f)

30 wt% and (g) 12 wt%. TiO2 coated CNTs in the presence of BA with the CNT

concentration of (h) 30 wt% and (i) 12 wt%. The arrows indicated the uncoated

(004) (100)

(002)

(j)

20 30 40 50 60

0

5000

10000

15000

20000

Inte

nsity / a

.u.

2θ / º

C (002)

(k)

Fmoc-Tyr

BA

(f) (g)

(h) (i)

201

part of CNTs. Note. This was different from the cracks resulting from the drying

effect. (j) SAED pattern taken from the sample shown in (f). (k) XRD pattern of the

as-produced CNT-TiO2 nanohybrids. C: CNT. For (c), (e), (g) and (i), scale bar =

200 nm. For (a), (b), (d), (f) and (h), scale bar = 100 nm.

Table 5.3 Measured thickness of the TiO2 coating based on the TEM observation

30 wt%

12 wt%

Fmoc-Trp 10.4 ± 2.3 48.3 ± 7.3

Fmoc-His 13.4 ± 1.1 49.7 ± 7.9

Fmoc-Tyr 12.7 ± 3.8 70.0 ± 2.2

BA 4.3 ± 1.0 17.7 ± 2.1

5.3.4 Effect of modifier to CNT ratio on the hybrid

morphology

Zhang et al. 16 highlighted the critical role of the ratio of surfactants to CNTs in

determining the morphology of the SiO2 coating. Herein, to investigate the effect

of the ratio of the Fmoc-AAs to CNTs on the morphology of the resulting hybrids,

Fmoc-His solution was diluted by a factor of 10 while the other parameters were

kept constant. SEM images revealed the strong influence of the concentration of

Fmoc-His on the morphology of the hybrids which undergoes a dramatic change

from uniformly coated CNTs at high Fmoc-His loadings (Figure 5.9a) to random

coexistence of isolated TiO2 particles and uncoated CNTs at low Fmoc-His

loadings (Figure 5.9b). The lack of coating at low Fmoc-His loading is believed

to be due to an insufficient density of Fmoc-His on CNTs’ surface to induce TiO2

deposition. This observation further emphasizes the role of the Fmoc-AAs in the

synthesis of the hybrid nanostructures.

Surface

modifier

Concentration

of CNTs

Thickness of

coating (nm)

202

Figure 5.9 SEM images of the structures produced with the addition of (a)

undiluted and (b) diluted Fmoc-His solutions (by a factor of 10).

5.3.5 Kinetics for TiO2 growth

In order to study the kinetics for the growth of TiO2 coating on Fmoc-AA

functionalized CNTs, the coating process was followed by taking SEM images of

the hybrids which had be coated for different times. Figure 5.10 shows the SEM

images taken after reaction for 10 min, 1 h and 6.5 h respectively. Due to the

similar issue as with the study of the kinetics for silica growth, herein we

evaluated the kinetics for TiO2 growth by considering the change in the diameter

of the hybrid tubes. The diameter of TiO2 coated CNTs was found to increase

from 68.0 ± 17.3 nm after growth for 10 min (Figure 5.10a) to 82.4 ± 21.5 nm

after growth for 1 h (Figure 5.10b) and to 94.2 ± 16.8 nm after growth for 6.5 h

(Figure 5.10c). Little change in the diameter of the hybrid NT was observed after

1h of reaction (Figure 5.10d), suggesting that the growth of TiO2 coating on

CNTs was saturated. It is worth noting that for the study of the growth of silica,

the reaction was still ongoing over a period of days.

(b)

1μm

(a)

203

Figure 5.10 SEM images of TiO2 coating growing on Fmoc-Trp functionalized

CNTs at different reaction times of (a) 10 min (b) 1 h and (c) 6.5 h. (d) Plot of the


calculated based on 50 separate measurements.

5.3.6 Synthesis of TiO2 NTs In order to induce crystallization of the coating and to produce TiO2 NTs, the

as-produced hybrid materials were calcined at 400 °C 17 for 2 h followed by at

550 ºC for 2 h to remove CNT templates. Previous studies have shown that 550ºC

was sufficiently high to oxidize CNT templates but not so high as to destroy the

TiO2 nanotube morphology 18. The oxidation proceeds from inside the hollow

region of CNTs rather than from outside graphitic layers 19.

The morphology of the thermally treated hybrids was characterized by SEM and

(a) (b)

(c)

0 1 2 3 4 5 6 7 8

0

50

100

150

Dia

me

ter

of th

e h

yb

rid

NT

/ n

m

Growth time / h

(d)

204

TEM. The SEM images shown in Figure 5.11 revealed the nanotubular structures

with a rather smooth surface indicating the small crystal size. The length of the

synthesized TiO2 NTs ranged from a few hundred nanometers to several microns

which was much shorter than that of the initial CNT templates. This was limited

by the fractured TiO2 coatings prior to calcination, thus indicating the effect of

drying process on the length of the synthesized NTs. The NTs with both open and

closed ends were observed as indicated by the arrows in Figure 5.11b. The

removal of CNT templates after calcination was confirmed by EDX. As shown in

Figure 5.11c, the three most distinct peaks were O, Ti and Al which were

originated from the TiO2 NTs and the SEM substrate respectively. The presence

of significantly weaker C signal compared with that of the as-produced hybrids

indicated that the CNT templates were largely removed after calcination at 550ºC.

Figure 5.11 SEM images of TiO2 nanotubes produced from (a) TiO2 coated

Fmoc-His functionalized CNTs (30 wt%) and (b) TiO2 coated Fmoc-Tyr

functionalized CNTs (12 wt%). (c) EDX spectrum of the hybrid after calcination at

(a) (b)

0 2 4 6

0

500

1000

1500

2000

2500

Co

un

ts (

a.u

.)

Energy (keV)

O

Al

Pt Ti

(c)

205

550 ºC. Note. Pt signal was originated from the conductive coating on the SEM

sample. Scale bar, (a) 500nm, (b) 1μm.

The TEM images gave further information about the structure of the produced

TiO2 NTs. As shown in Figure 5.12, calcination of the hybrids produced with

30wt% of CNTs resulted in very thin and collapsed NT structures (Figure 5.12a,e

and g) except for that produced from Fmoc-His functionalized CNTs (Figure

5.12c) which exhibited well-defined NT structures with the inner diameter of

58.9± 17.4 nm and wall thickness of 12.5 ± 1.3 nm. TiO2 NTs with uniform inner

diameter and wall thickness were obtained after calcination of the hybrids

produced with 12wt% of CNTs (Figure 5.12b,d and f) except for that produced in

the presence of BA which showed a collapsed tubular structure (Figure 5.12h).

The measured inner diameter and wall thickness of all the synthesized TiO2 NTs

were summarized in Table 5.4 and 5.5 respectively.

Table 5.4 Measured inner diameter of the synthesized TiO2 NTs

30wt%

12wt%

Fmoc-Trp very thin and collapsed NTs 62.5 ± 13.0

Fmoc-His 58.9 ± 17.4 50.6 ± 17.4

Fmoc-Tyr very thin and collapsed NTs 56.2 ±12.6

BA collapsed NTs collapsed NTs

Table 5.5 Measured wall thickness of the synthesized TiO2 NTs

30wt%

12wt%

Fmoc-Trp very thin and collapsed NTs 33.0 ± 6.0

Fmoc-His 12.5 ± 1.3 23.9 ± 3.2

Fmoc-Tyr very thin and collapsed NTs 33.6 ±7.3

BA collapsed NTs collapsed NTs

The mesoporous nature of the synthesized TiO2 NTs was evidenced by the high

range in contrast in the micrograph. Higher magnification image revealed that the

TiO2 NT was composed of many spherical nanocrystals with 9.9 ± 1.6 nm in

Surface

modifier

Inner diameter

of TiO2 NT (nm)

Concentration

of CNTs

Surface

modifier

Concentration

of CNTs

Wall thickness

of TiO2 NT (nm)

206

diameter. Only isolated NTs were obtained without the presence of

interconnected NT network, indicating that the CNTs were well dispersed in the

presence of the Fmoc-AAs. Both open and closed end NTs were observed as

indicated by the arrows in Figure 5.12b and Figure 5.12f which was consistent

with the SEM observation. It was also noted that the inner diameter of the

synthesized TiO2 NTs is almost consistent with the outer diameter of the

carbonaceous template, indicating that both the inner diameter and wall thickness

of the TiO2 NTs can be controlled by varying the dimension of CNT templates

and the CNT to TBOT ratio respectively.

The crystal structure and phase composition of the produced TiO2 NTs were

examined by SAED. As shown in Figure 5.12(i-l), the SAED patterns taken from

the calcined samples (upper half) showed excellent agreement with those of

anatase as simulated from JCPDS 21-1272 (lower half) indicating the

crystallization of amorphous TiO2 into anatase. Besides, the observed ring

patterns instead of diffraction spots suggest the polycrystalline nature of the

anatase NTs as a result of the random orientation of the TiO2 nanocrystallites.

The absence of the diffraction corresponding to the graphitic structure confirms

the complete removal of the carbonaceous template after calcination which agrees

well with the EDX result. XRD analysis was also conducted to confirm the

anatase phase of the TiO2 NTs. As shown in Figure 5.12m, upon calcination at

400 ºC followed by at 550 ºC, the amorphous coating crystallized into anatase

and the carbonaceous template was completely removed as evidenced by the

presence of the diffraction peaks at 2θ=25.4º, 37.9º, 48.1°, 54.0° and 55.2° which

are typical for anatase phase (JCPDS 84-1286) and the absence of the carbon

(002) diffraction, respectively. This result is consistent with the SAED study. The

considerably lower oxidation temperature for the CNT templates compared with

that for the pristine CNTs which was ~650-700 ºC 2,20,21 may be attributed to the

catalyzing effect of TiO2 coating on the carbon gasification via the Mars-van

Krevelen mechanism 2,18. This observation is in contrast to that observed for

CNT-SiO2 hybrids which exhibited a higher oxidation resistance 22. This may be

207

explained by that the presence of SiO2 coating tended to hinder the thermal

decomposition of CNTs by preventing the access of oxygen 23.

The size of the anatase nanocrystallites was estimated from the XRD pattern

using Scherrer equation, following:

cos

Kd (5.1)

Where, d is the crystal size of anatase, K is the shape factor (usually taken as 0.9),

λ is the wavelength of X-ray radiation, β is the full width of the anatase (101)

peak at half maximum intensity (FWHM, in radians) and θ is the corresponding

diffraction angle. The peak was fitted by Gaussian function. The value of β was

determined using OriginPro 8.1 software. The average crystal size of anatase was

calculated to be ~14 nm which was comparable to that reported by Eder 2. The

slightly higher value of the crystal size measured according to XRD compared

with that determined by TEM could be attributed to the Scherrer equation tending

to overestimate the size of particles with diameters less than 10 nm 24. It should

be also noted that the factors such as the instrumental broadening effect and strain

(lattice distortion) were not considered in the present study which also contribute

to the peak broadening.

208

(a) (b)

(g)

30 wt% CNTs

12 wt% CNTs

Fmoc-Trp

Fmoc-His

(c) (d)

(e) (f)

(h)

Fmoc-Tyr

BA

209

20 30 40 50 60

0

10000

20000

30000

40000

Inte

nsity / a

.u.

2θ / º

A (101)

A (004) A (200)

A (105) A (211)

(m)

(101)

(004)

(200)

(211)

(101)

(004)

(200)

(211)

(i)

(101)

(004)

(200)

(211)

(k)

(101)

(004)

(211)

(200)

(l)

(101)

(004)

(200)

(211)

(j)

210

Figure 5.12 TEM images of the calcined hybrids. (a) In the presence of Fmoc-Trp

and 30wt% of CNTs. (b) In the presence of Fmoc-Trp and 12wt% of CNTs. (c) In

the presence of Fmoc-His and 30wt% of CNTs. (d) In the presence of Fmoc-His and

12wt% of CNTs. (e) In the presence of Fmoc-Tyr and 30wt% of CNTs. (f) In the

presence of Fmoc-Tyr and 12wt% of CNTs. (g) In the presence of BA and 30wt% of

CNTs. (h) In the presence of BA and 12wt% of CNTs. (i-l) SAED patterns taken

from the samples shown in (b-d) and (f) respectively (upper half) which confirmed

the polycrystalline anatase phase of the NTs by showing excellent agreement with

those simulated from JCPDS 21-1272 (lower half). The SAED patterns were

indexed to the (101), (004), (200) and (211) planes of anatase phase. (m) XRD

pattern taken from the sample shown in (d). A: anatase. For (a), (e) and (g), scale

bar = 100 nm and for (b), (c), (d), (f) and (h), scale bar = 200 nm.

The lattice fringes with an interplanar d-spacing of 0.35 nm was clearly visible in

the HRTEM image (Figure 5.13) which correspond to the (101) planes of anatase

(JCPDS 21-1272). The HRTEM image also reveals the different orientations of

the (101) crystal planes, suggesting that there was no specific orientation of the

TiO2 nanocrystallites. This observation is consistent with the ED patterns.

0.35 nm

A (101)

0.35 nm

A (101)

211

Figure 5.13 HRTEM image of a synthesized TiO2 NT showing the lattice spacing of

0.35 nm, corresponding to the (101) crystal planes of anatase. Scale bar, 10 nm.

5.3.7 Phase transformation

The transition from anatase to rutile is kinetically unfavorable at lower

temperatures. Therefore, temperatures between 600 ºC and 900 ºC is normally

required to induce phase transformation 25. However, at such temperatures, the

unsupported TiO2 NTs tend to collapse due to the stress associated with the

reconstruction of phase. Eder et al. 18 have shown that CNTs can support the

anatase coating during the phase transformation and be subsequently removed to

produce phase-pure rutile nanotubes.

Herein, the as-produced hybrids were heated in argon to prevent CNT oxidation.

An additional benefit of argon is that it has been previously reported that heating

in argon can lower the phase transformation temperature by at least 100 ºC 2. The

influence of heating temperature, pre-treatment and ramp rate on the phase

transformation was investigated. First, to study the effect of heating temperature,

the as-produced hybrids were heated in argon at 900 ºC and 800 ºC for 2 h using a

ramp rate of 20 ºC/min for phase transformation. Both samples were then

calcined at 550 ºC with the aim of removing the CNT templates. As shown in

Figure 5.14a, the CNTs survived the heat treatment in argon at 900 ºC followed

by in air at 550 ºC. Rod-like TiO2 nanocrystals with an average length of ~ 74 nm

and a diameter of ~48 nm were deposited on the surface of the nanotubes. These

nanocrystals were aligned approximately parallel to the long axis of the CNTs

and the XRD analysis showed that they were pure rutile phase (Figure 5.14b).

The average size of the rutile crystals was also determined from the X-ray line

broadening to be ~ 20 nm. The large discrepancy between the value calculated

from the XRD data and that from the TEM observation may due to the fact that

the shape factor K of 0.9 is only used for spherical particles, so it gave a poor

estimation of the particle size for rod-like crystallites. In contrast, the CNTs were

burnt out to leave anatase NTs after heating the as-produced hybrids in argon at

800 °C followed by in air at 550 ºC (Figure 5.14c) as supported by the SAED

212

pattern (Figure 5.14d, upper half) which showed excellent agreement with those

of anatase as simulated from JCPDS 21-1272 (lower half). This observation is in

contrast to the study which demonstrated that the onset of transition from anatase

to rutile was above 700 °C 2,18.

Furthermore, two different ways of inducing phase transformation were compared:

(1) to heat the as-produced hybrids directly in argon at 800 ºC and (2) to heat the

hybrids first in air at 400 ºC to induce crystallization followed by heating the

anatase coated CNTs in argon at 800 ºC. The ramp rates were kept at 20 ºC/min

for both methods. In contrast to those shown in Figure 5.14c, TEM image of the

sample prepared using the 2nd method showed that spherical TiO2 nanocrystals

with the average size of ~43 nm were deposited on CNTs’ surface (Figure 5.14e).

Corresponding XRD pattern (Figure 5.14f) revealed the presence of new

diffraction peaks typical for rutile phase (2θ=27.5°,36.1°,41.3°,44.1°, 54.4° and

56.7°, JCPDS 88-1175) in addition to anatase, indicating partial phase

transformation. Similar observation has previously been reported by Eder 3 who

demonstrated that the presence of BA tended to retard the phase transformation

which was attributed to the increased surface strain induced by the adsorbed

surfactants. The size of the anatase and rutile crystals was calculated from the

XRD data to be ~32 nm and ~40 nm respectively which showed moderate

agreement with the TEM study. The mass fraction of rutile phase which indicated

the extent of phase transformation was also determined from the relative

integrated peak intensities, following3:

R

A

R

I

IW

88.01

1

(5.2)

Where, WR is the percentage of rutile, IA and IR are the integrated intensity of

anatase (101) and rutile (110) peaks respectively, which were determined using

OriginPro 8.1 software. The calculated value indicated that more than 60% of

anatase has transformed to rutile. Mehranpour et al. 26 have observed complete

213

phase transformation at 800 ºC. The slower kinetics of phase transformation in

the current study may due to the non-uniform morphology and wide size

distribution of the TiO2 nanoparticles as well as the different method of

preparation 27.

Finally, the effect of ramp rate was studied with the heat treatment performed at

20 ºC/min and 1 °C/min respectively. With 1 ºC/min of heating rate, closely

packed TiO2 nanocrystals with a much smaller size of ~17 nm were observed to

uniformly coat CNTs (Figure 5.14g), indicating that the lower heating rate

favored the formation of smaller crystals and denser coverage. SAED analysis

(Figure 5.14h, upper half) confirmed the rutile phase of the TiO2 nanocrystals by

showing agreement with the diffractions of rutile as simulated from JCPDS

21-1276 (Figure 5.14h, lower half).

(c)

20 30 40 50 60

0

1000

2000

3000

4000

5000

6000

Inte

nsity / a

.u.

2θ /

R

R

R R

R

(b)

2θ / °

(d)

(211)

(200)

(004)

(101)

(a)

214

Figure 5.14 (a) TEM images of the hybrids after heat treatment in Ar at 900 ºC followed by

in air at 550 ºC with the ramp rate of 20 ºC/min. Scale bar, 20 nm. (b) XRD pattern taken

from the sample shown in (a). (c) TEM image of the hybrids after heat treatment in Ar at

800 ºC followed by in air at 550 ºC with the ramp rate of 20 ºC/min. Scale bar, 100 nm. (d)

SAED pattern (upper half) taken from the sample shown in (c). The pattern was indexed to

the (101), (004), (200) and (211) planes of anatase phase. (e) TEM image of the hybrids after

heat treatment in air at 400 ºC followed by in Ar at 800 ºC with a ramp rate of 20 ºC/min.

Scale bar, 100 nm. (f) XRD pattern taken from the sample shown in (e). A: anatase, R: rutile,

C: CNT. (g) TEM image of the hybrids after heat treatment in air at 400 ºC followed by in

Ar at 800 ºC with a ramp rate of 1 ºC/min. Scale bar, 200 nm. (h) SAED pattern (upper half)

taken from the sample shown in (g). The SAED pattern was indexed to the (110), (111), (210),

(211) and (220) planes of rutile phase. All the heat treatments were conducted for 2h.

5.3.8 Aligned arrays of TiO2 NTs

Vertically aligned CNT arrays with the length of up to ~400 μm and area of 5 mm

(f)

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(220)

(211) (h)

215

x 5 mm were used for the templating of TiO2 NT arrays on Si/SiO2 substrate.

SEM analysis of the sample produced in the absence of the surfactants showed

randomly oriented nanofibers between which sheets had grown (Figure 5.15a).

This observation could be explained by that in the presence of pristine CNTs,

there was no driving force for the diffusion of TBOT into the hydrophobic arrays,

which instead formed a film on the top of the arrays. Moderately aligned

nanofibrous structures were obtained in the presence of Fmoc-Trp (Figure 5.15b),

Fmoc-His (Figure 5.15c) and Fmoc-Tyr (Figure 5.15d). The length of the

resultant arrays was up to ~300 μm, indicating that the surfactants could penetrate

into the arrays to adsorb on the lower portion of CNTs and subsequently promote

the diffusion of TBOT. It was noted in Figure 5.15(b) and (d) that although the

TiO2 NTs were well aligned within the arrays, the vertical alignment of the arrays

with respect to the substrate was lost, possibly during (1) the adsorption of

Fmoc-AAs in solution or (2) the coating process in TBOT/EtOH solution or (3)

the calcination process. Whilst the bundled up arrays shown in Figure 5.15c

revealed a reasonably good vertical alignment with the upper part decurved. In

contrast, the products obtained in the presence of BA showed broken NT

structures due to the low density of deposited TiO2 nanoparticles on CNTs prior

to calcination (Figure 5.15e). This result further supported that the Fmoc-AAs

were more efficient binders for CNTs than BA.

216


where as-produced CNT mat was used as templates. TiO2 NT arrays produced in

the presence of (b) Fmoc-Trp (c) Fmoc-His (d) Fmoc-Tyr and (e) BA.

TEM analysis was also conducted to further characterize the internal structures of

the products. In contrast to the control experiment which resulted in the formation

of collapsed TiO2 NTs (Figure 5.16a), well-defined TiO2 NTs with relatively

(a)

(c) (d)

(b)

(e)

217

uniform inner diameter and wall thickness were obtained in the presence of the

Fmoc-AAs (Figure 5.16b-d). The measured inner diameter and wall thickness of

the resultant NTs were summarized in Table 5.6. It was noted that the inner

diameter of the produced TiO2 NTs was smaller than the outer diameter of the

CNT templates (56.6 ± 16.6 nm) which could be attributed to the stress induced

contraction during the oxidative removal of CNTs. Black arrows in (c) indicated

the presence of residual CNT templates after calcination. In the presence of BA,

collapsed NT structures were observed (Figure 5.16e) due to the weaker π–π

stacking interactions between BA and CNTs. The TEM result agreed well with

the SEM observation. XRD analysis confirmed the anatase phase of the TiO2 NTs.

The absence of carbon (002) diffraction indicated the complete removal of CNT

templates after calcination.

(a) (b)

(c) (d)

218

Figure 5.16 TEM images of (a) the product obtained from the control experiment.

TiO2 NTs produced in the presence of (b) Fmoc-Trp (c) Fmoc-His and (d)

Fmoc-Tyr. (e) Collapsed NT structures obtained in the presence of BA. The red

arrow in (b) and (c) indicated the open ends of the TiO2 NTs. Note. CNT templates

were not completely removed after calcination as indicated by the black arrows in

(c). (f) XRD pattern taken from the sample shown in (b). Scale bar, (a-e) 200 nm.

Table 5.6 Measured inner diameter and wall thickness of the resultant TiO2 NTs

Inner diameter (nm) Wall thickness (nm)

Fmoc-Trp 44.3 ± 10.0 9.5 ± 2.2

Fmoc-His 50.0 ± 15.1 10.7 ± 1.7

Fmoc-Tyr 38.0 ± 10.1 7.7 ± 0.7

5.4 Combined sites for catalyzing SiO2 and TiO2

deposition

Morse et al. 11 demonstrated the synthesis of silica catalyzed by the interface of

two populations of gold nanoparticles, one of which had either been

functionalized with serine residues and the other with histidine residues. This

interface between the two residues mimicked the catalytically active site of

silicateins. This concept was developed further in this study, by combining two

populations of CNTs, one functionalized with Fmoc-His and the other with

Fmoc-Tyr which expose imidazole (hydrogen bonding) and hydroxyl

(nucleophilic) functionalities respectively. These combined populations were then

studied for the biomimetic catalytic synthesis of silica and titania.

(e)

20 30 40 50 60

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6000

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nsity / a

.u.

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219

5.4.1 Synthesis of the biomimetic catalyst

The assembly of two populations of functionalized CNTs was initially studied.

SEM characterization revealed two distinct morphologies; bundled fibers (Figure

5.17a) and spherical aggregates with an average diameter of 37.7 ± 8.0 μm

(Figure 5.17b). High magnification micrographs (Figure 5.17c and d) revealed

densely entangled CNTs in both aggregates. The driving force for this

aggregation was believed to be the H-bonding between the two populations of

functionalized CNTs. A control experiment was also conducted where CNTs

functionalized with single Fmoc-AA species was characterized by SEM. Both

Fmoc-His functionalized CNTs (Figure 5.17e) and Fmoc-Tyr functionalized

CNTs (Figure 5.17f) showed only randomly aligned CNT networks.

(a)

(d) (c)

(b)

220

Figure 5.17 SEM images showing (a) bundled fibers and (b) spherical aggregates

formed in the combined solutions. (c,d) Magnified images of the aggregates shown

in (a) and (b) respectively. (e) Fmoc-His f-CNTs and (f) Fmoc-Tyr f-CNTs.

5.4.2 Synthesis of SiO2 catalyzed by the combined sites

The templating of the combined catalyst system was subsequently studied. The

combined populations of nanotubes were first deposited on a silicon wafer, which

was then immersed in a TEOS solution for 6 d. The product was then annealed in

air at 200 ºC for 3 h followed by annealing at 650 ºC for 2 h to remove the

carbonaceous template. The thermally treated sample comprised of ribbons

approximately 100-500 nm wide surrounded by short fibres (Figure 5.18a and b).

These ribbons are believed to be the result of coating the bundled fibres of CNTs.

However, no structures relating to the spherical aggregates were observed. EDX

analysis (Figure 5.18c) confirmed the silica nature of the annealed sample and the

presence of rather weak C signal indicated that the CNT templates were largely

removed after calcination.

(e) (f)

221

Figure 5.18 (a,b) SEM image of silica coated combined catalyst after heat treatment.

(c) EDX spectrum of the sample shown in (a).

5.4.3 Synthesis of TiO2 catalyzed by the combined

sites

The silicon wafer coated with the dried combined populations of nanotubes was

also immersed into a TBOT solution. TiO2 nanorods with an average length of

114.4 ± 26.4 nm and width of 31.0 ± 5.4 nm were uniformly deposited on both

bundled NTs (Figure 5.19a) and individual NTs (Figure 5.19b). EDX analysis

(a)

0 2 4 6

0

5000

Inte

nsity

Energy (keV)

C

O Cu

Si

Al

(c)

(b)

222

shown in Figure 5.19e confirms the TiO2 coating on CNT bundles. Such coating

morphology is different from those observed with the reaction proceed in solution

which led to rather smooth TiO2 coating (section 5.3.1). After heat treatment in

air at 400 ºC followed by at 550 ºC, TiO2 particle aggregates as well as thin fibres

were observed (Figure 5.19c). EDX analysis confirms the removal of CNT

templates after calcination by showing a negligible C signal (Figure 5.19f). A

control experiment with direct depositing TBOT solution on Si wafer was

performed (Figure 5.19d). TiO2 nanorods with similar dimensions to those

obtained in the presence of the combined sites were observed to coexist with

larger TiO2 particles. The exact mechanism for the formation of TiO2 nanorods is

unclear at this stage.

(a

)

(c

) (d)

(b)

223

(e) (f)

Figure 5.19 SEM images of (a) TiO2 nanorods coated CNT bundles (b) TiO2

nanorods coated individual CNTs (c) TiO2 nanorods coated CNT bundles after heat

treatment and (d) TiO2 particles formed on Si wafer. (e) and (f) EDX spectrum of

the sample shown in (a) and (c) respectively.

5.5 Conclusion

In this chapter, an in-situ sol-gel method was demonstrated for the synthesis of

CNT-SiO2 and CNT-TiO2 nanohybrids at low temperature and neutral pH. Both

SEM and TEM observation confirm the formation of uniform SiO2 coating from

TEOS on individual CNTs in the presence of Fmoc-Trp and Fmoc-His, indicating

good dispersion of CNTs. Fmoc-Tyr functionalized CNTs resulted in a mixture of

partially coated and uncoated CNTs due to the poorer affinity of Fmoc-Tyr for

CNTs as well as its weaker H-bonding interaction with the hydrolyzed TEOS. A

similar dependence of the degree of coating on the Fmoc-AA structure was

observed for coating CNTs with TiO2 from TBOT, except for the case of

Fmoc-Tyr which led to uniform coating of TiO2. The morphology of the resultant

hybrids was found to be highly dependent on the CNT to TBOT ratio and the

modifier to CNT ratio. The coating of CNTs functionalized with BA, as reported

by Eder et al. was also studied. It was found that the TiO2 coating obtained on

these BA functionalized CNTs was not as uniform as on the Fmoc-AA

0 2 4 6

0

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400

600

800

1000

1200

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.)

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1500

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.)

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224

functionalized nanotubes. This difference in coating was attributed to the weaker

π–π stacking interactions between BA and CNTs as a result of the lower degree

of aromaticity of the benzyl ring.

The formation mechanism of the SiO2 and TiO2 coating on the Fmoc-AAs

functionalized CNTs was proposed. The surface modifiers were believed to play a

dual role: (1) To stabilize CNT dispersion by acting as electrostatic surfactants. (2)

The presence of indole, imidazole and hydroxyl functionalities from the side

chains of the amino acid moieties provide binding sites for SiO2 and TiO2

deposition.

Anatase TiO2 NTs were synthesized after calcination of the CNT-TiO2

nanohybrids as confirmed by XRD and ED analysis. HRTEM image revealed that

these NTs were the aggregates of many TiO2 nanocrystallites with no preferential

orientation. Both the inner diameter and wall thickness of the synthesized TiO2

NTs can be controlled by varying the dimension of CNT templates and the ratio

of CNT to TBOT. The transition from the anatase to rutile phase was achieved via

heat treatment of the hybrids in argon to prevent CNT oxidation. Heating

temperature, pre-treatment and ramp rate were found to affect the phase

transformation.

A simple route toward the production of TiO2 NT arrays was also demonstrated

using the CVD grown vertically aligned CNT arrays as templates in the presence

of the Fmoc-AAs. The length of the TiO2 NT arrays was up to ~300 μm,

benefiting from the avoidance of extended sonication. SEM analysis confirmed

the good alignment of TiO2 NTs in the arrays although the vertical alignment

with respect to the substrate was lost. Such structures may be useful for the

development of electronic units.

The noncovalent approach presented in this study permits the preservation of the

structure and properties of pristine CNTs. The template-directed synthesis

225

technique also allows control over the morphology of the inorganic

nanostructures.

5.6 References

1. L. Zhao et al., Coating of multi-walled carbon nanotubes with thick layers of

tin(IV) oxide, Carbon, 2004, 42, 1858.

2. D. Eder et al., Morphology control of CNT-TiO2 hybrid materials and rutile

nanotubes, J.Mater.Chem., 2008, 18, 2036.

3. D. Eder et al., Carbon–Inorganic Hybrid Materials: The Carbon-Nanotube/

TiO2 Interface, Adv.Mater., 2008, 20, 1787.

4. M. Liang et al., Imidazole catalyzed silica synthesis: Progress toward

understanding the role of histidine in (bio)silicification, J. Mater. Res., 2009, 24,

1700.

5. S. V. Patwardhan et al., Silicification and Biosilicification. Part 6.

Poly-L-Histidine Mediated Synthesis of Silica at Neutral pH,

J.Inorg.Organomet.Polym., 2003, 13, 49.

6. F. Rodríguez et al., Study of the Chemical and Physical Influences upon in

Vitro Peptide-Mediated Silica Formation, Biomacromolecules, 2004, 5, 261.


highly ordered TiO2 nanotube arrays, Environ. Sci. Technol., 2008, 42, 8547.

8. G. K. Mor et al., Use of Highly-Ordered TiO2 Nanotube Arrays in

Dye-Sensitized Solar Cells, Nano Lett., 2006, 6, 215.

9. M. S. Sander et al., Template-Assisted Fabrication of Dense, Aligned Arrays of

Titania Nanotubes with Well-Controlled Dimensions on Substrates, Adv. Mater.,

2004, 16, 2052.



11. D. Kisailus et al., Functionalized Gold Nanoparticles Mimic Catalytic Ativity

of a Polusiloxane-Synthesizing Enzyme, Adv. Mater., 2005,17,1234.

12. J.N.Cha et al., Biomimetic synthesis of ordered silica structures mediated by

block copolypeptides, Nature, 2000, 403, 289.

13. Y. Zhou et al., Efficient Catalysis of Polysiloxane Synthesis by Silicatein a

Requires Specific Hydroxy and Imidazole Functionalities, Angew. Chem. Int. Ed.,

1999, 38, 779.

14. L.W. Kelts et al., Sol-gel chemistry studied by 1H and 29Si nuclear magnetic

resonance, Journal of Non-Crystalline Solids, 1986, 83, 353.

15. J. K. Park et al., J. Korean Chem. Soc. 1998, 22, 281.

16. M. Zhang et al., Fabrication of mesoporous silica-coated CNTs and

application in size-selective protein separation, J. Mater. Chem., 2010, 20, 5835.

17. Z. Németh et al., Preparation of homogeneous titania coating on the surface

of MWNT, Composites Science and Technology, 2011, 71, 87.

18. D. Eder et al., Pure rutile nanotubes, Chem. Commun., 2006, 1448.

19. B. C. Satishkumar et al., Oxide nanotubes prepared using carbon nanotubes as

http://link.springer.com/search?facet-author=%22Siddharth+V.+Patwardhan%22

226

templates, J.Mater.Res., 1997, 12, 604.

20. O. M. Yaghi et al., Selective binding and removal of guests in a microporous

metal–organic framework, Nature, 1995, 378, 703.

21. P. M. Ajayan et al., Carbon nanotubes as removable templates for metal oxide

nanocomposites and nanostructures, Nature, 1995, 375, 564.


Chem.Commun., 2002, 34.

23. A. B. Bourlinos et al., Preparation of a water-dispersible carbon

nanotube–silica hybrid, Carbon, 2007, 45, 2136.

24. S. P. Sree et al., Synthesis of uniformly dispersed anatase nanoparticles inside

mesoporous silica thin films via controlled breakup and crystallization of

amorphous TiO2 deposited using atomic layer deposition, Nanoscale, 2013, 5,

5001.

25. S. R. Yoganarasimhan et al., Mechanism of crystal structure transformations.

Part 3.—Factors affecting the anatase-rutile transformation, Trans. Faraday Soc.,

1962, 58, 1579.

26. H. Mehranpour et al., Study on the Phase Transformation Kinetics of

Sol-Gel Drived TiO2 Nanoparticles, Journal of Nanomaterials, Volume 2010,

Article ID 626978, doi:10.1155/2010/626978.

27. Y. Iida et al., Grain Growth and Phase Transformation of Titanium Oxide

During Calcination, J. Am. Ceram. Soc., 1961, 44, 120.

227

Chapter 6 Mineralization of peptide self-assembled

hydrogels

6.1 Introduction

Xu et al. have reported the self-assembly of Fmoc-AA hydrogelator induced by

enzyme 1. They have proposed a model for the supramolecular structure of the gel

based on the spectroscopic studies which suggested the presence of large number

of hydroxyl groups on the nanofibres’ surface. The presence of such high density

of -OH groups may catalyze the deposition of silica. Therefore, in the current

study, enzyme triggered Fmoc-Y and Fmoc-FY self-assembled hydrogels were

employed as templates for silica deposition.

6.2 Enzymatic self-assembly of Fmoc-Y and Fmoc-FY

hydrogels

6.2.1 Fmoc-Y hydrogel

Alkaline phosphatase (AP)-catalyzed dephosphorylation convert the ionic PO43-

group of Fmoc-tyrosine (phosphate)-OH (Fmoc-Y(p)-OH) into a neutrual group

(Figure 6.1a), leading to a hydrogelator Fmoc-Tyr (Fmoc-Y) that self-assembled

into supramolecular hydrogels. Negatively stained TEM image of the diluted

hydrogel revealed a network of nanofibres with a uniform width of 19.3 ± 3.2 nm

(Figure 6.1b). The gel appeared opaque indicating the presence of wider fibres or

bundles of self-assembled nanofibrils 1. The undiluted hydrogel showed a mixture

of nanofibrils and nanoribbons (Figure 6.1c and d). The ribbons were quite flat

while the nanofibrils revealed a very regular morphology with distance between 2

successive ‘pinched joints’ of 270.8 ± 21.9 nm (and width at the widest section

being 33.2 ± 4.0 nm). Such twisted morphology was not observed in previous

studies.

228

Error!

Figure 6.1 (a) Schematic representation of the enzymatic dephosphorylation of

Fmoc-Y(p)-OH to Fmoc-Y. The corresponding optical images for Fmoc-Y(p)-OH

precursor solution before enzyme addition and the self-supporting hydrogels

formed were also shown. (b) Negatively stained TEM image of the diluted Fmoc-Y

hydrogel. Scale bar, 50 nm. (c,d) Negatively stained TEM image of the undiluted

hydrogel. Scale bar, 100 nm.

nanoribbon

nanofibres

(d)

Alkaline Phosphatase

pH 7.03, 37ºC

(a)

Fmoc-Y(p)-OH Fmoc-Y

(b

)

270nm

100 nm

(c)

229

6.2.2 Fmoc-FY hydrogel

The 2nd template studied is Fmoc-FY self-assembly hydrogel. The self-assembly

mechanism for Fmoc-FY is better established and the Fmoc-FY hydrogel is more

stable and reproducible than Fmoc-Y hydrogel. AP catalyzed dephosphorylation

converts the ionic PO43- group of Fmoc-FpY into a neutrual –OH group, leading

to the transformation from a micellar solution into a fibrous hydrogel2 (Figure

6.2a). The phosphorylated hydroxyl group of Fmoc-FpY adds a polar character to

the C-terminus, a micellar structure is therefore expected to form with the

hydrophobic Fmoc group sequestered inside and the hydrophilic phosphate head

group in contact with the solvents.

TEM image showed nanofibres with the average width of 24.0 ± 2.3 nm (Figure

6.2b). HPLC trace was recorded to monitor the dephosphorylation process. As

shown in Figure 6.2c, the gelation occurred when ~70% Fmoc-FpY have been

converted which took ~ 90 min. The supramolecular arrangement within both the

micelles and fibres was further studied by fluorescence emission spectroscopy.

The solution of Fmoc-FpY exhibited a peak centered at 320 nm (Figure 6.2d)

which was attributed to the unassembled Fmoc group and a shoulder between 350

and 370 nm. Upon addition of AP, the Fmoc peak red shifted slightly to 325 nm,

whereas the shoulder peak disappeared. The loss of the shoulder peak at 370 nm

indicated the disassembly of the micelles 2. The decreased intensity of 320 nm

peak as well as a small red-shift indicated the self -assembly of Fmoc-FY into

fibre networks 3. This result is consistent with the previous study on the

production of Fmoc-FY gel.

230

(a)

(b)

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50

60

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80

0 20 40 60 80 100

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(c)

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Figure 6.2 (a) AP catalyzed dephosphorylation reaction of Fmoc-FpY and a

schematic representation of the supramolecular transition from micelles to fibres2.

(b) Negative stained TEM image showing the Fmoc-FY self-assembled nanofibrils.

Scale bar, 100 nm. (c) HPLC trace of the conversion of Fmoc-FpY to Fmoc- FY as a

function of time. The gelation point is marked with an arrow. (d) Fluorescence

emission spectra of the solution of Fmoc-FpY and the hydrogel of Fmoc-FY.

6.3 Silicification of the hydrogel nanostructures

6.3.1 Silicification of Fmoc-Y gel

6.3.1.1 Silicification via vortexing TEOS in the diluted

hydrogels (Method 1)

TEM was conducted to characterize the structure of the silicified nanostructures.

Two silicified structures were observed, i.e. coated nanofibres and coated

nanoribbons. The silica coating on the small nanofibres was not continuous, but

rather consisted of a very regular 25.2 ± 4.5 nm long coated regions with a

periodicity of 35 ± 5.6 nm (Figure 6.3a). The periodicity of the coating was not

observed on the nanoribbons. The periodicity of the silica coating was not

correlated with the helical pitch of the peptide nanofibres. Coatings with an

average thickness of ~ 8.4 ± 2.4 nm were observed on individual nanofibril after

reaction for 1 h with only a small number of non-templated silica present

0

50

100

150

200

250

300 350 400 450 500 550 600

Wavelength / nm

Inte

ns

ity

/ a

.u.

Fmoc-FpY solution

Fmoc-FY hydrogel

(d)

370 nm

320 nm

232

indicating the high affinity of the precursor for the peptide fibrils (Figure 6.3a).

After reaction for 2 h (Figure 6.3b) and 4 h (Figure 6.3d), the coating thickness

increased to 10.6 ± 3.5 nm and 10.4 ± 3.8 nm respectively. The coating of silica

on the fibrils was confirmed by EDX (Figure 6.3f).

The coating on the nanoribbons had a very different morphology to that on the

nanofibres. For the nanoribbons, the coating consisted of many aggregated

individual particles (~ 3.3 ± 0.7 nm in diameter, Figure 6.3b). The distribution of

the overall coating thickness was found to be broad and ranged from 13 to 48 nm.

One possibility is that the particles in these coatings were formed in solution by

the polymerization of TEOS and then deposited onto the ribbons. This will be

tested by using a control of TEOS vortexed in buffer solution to find if such

particles were produced.

One final observation which needs further investigation and confirmation is that it

appeared that the ratio of nanoribbons to nanofibres increased over reaction time.

Few ribbons were present during the early stage of the reaction (Figure 6.3a),

whilst large, broad ribbons were seen after 5 h (Figure 6.3e) which had formed

interconnected structures.

(a)

233

(b)

(c)

(d)

234

Figure 6.3 TEM images of silica coating on Fmoc-Y self-assembled nanostructures

after reaction for (a) 1 h, (b) 2 h, (c,d) 4 h and (e) 5 h. Scale bar, 100 nm. (f) EDX

spectrum of the mineralized peptide nanofibrils. (f) EDX spectrum of the silicified

fibrils.

6.3.1.2 Silicification via depositing TEOS/H2O mixture on the

hydrogels (Method 2)

0h

TEOS

H2O

10μm

(a)

(e)

(f)

Energy (keV)

Co

un

ts

6.05.04.03.02.01.00.0

30

20

10

0

Si

O

C

235

Figure 6.4 Silicification process of Fmoc-Y hydrogel. (a) SEM image of the dried

film of the hydrogel showing bundles of self-assembled nanofibres. (b) TEM

analysis indicated the spherical feature of the precipitate formed at the interface

with diamater of ~130 nm and (c) EDX spectrum confirmed the silica nature of the

precipitate.

A volume of Fmoc-Y hydrogel at 40 mM was covered with a mixture of

TEOS/H2O (volume ratio of hydrogel:TEOS:H2O=1:1:1). Phase separation was

observed immediately with TEOS phase present on the top due to its lower

density. SEM image of the dried gel (Figure 6.4a) showed the fibrillar structure.

Upon aging, a thin white layer was formed at the interface between the gel and

aqueous phase (indicated by arrow in Figure 6.4) and precipitation occurred in the

aqueous phase. TEM and EDX analysis on the white layer (Figure b and c)

revealed the formation of silica nanoparticles (Fig.). After aging for 1 month, a

clear region appeared in the upper part of the gel zone (indicated by arrow in

Figure 6.4), while the translucent hydrogel receded, suggesting the dilution of the

hydrogel.44 TEOS was completely consumed as evidenced by the disappearance

of the top phase, leaving only two phases separated by the silica layer. SEM

1 month

(b)

(c) 2 weeks

Clear region

White

layer

236

observation of the upper aqueous phase revealed that the hydrogel nanofibrils

were embedded in the silica nanoparticle aggregates (Figure 6.5a) as verified by

EDX (Figure 6.5b). It was noted that the hydrogel nanofibrils in the composite

appeared much thinner than those observed in the bottom gel phase (Figure 6.5c).

This may due to the dilution-induced disassembly of the thicker nanofibrils in the

aqueous phase. Interestingly, silica nanoparticles were also observed along with

the hydrogel fibrils in the bottom hydrogel phase (Figure 6.5c), which may be

introduced during sample collection (pipette tip breaks the silica layer and

pipetting silica together with the gel). A control experiment in which TEOS was

deposited on equal volume of H2O was also conducted. After standing for 2

month, no precipitation was observed and only the volume of TEOS reduced due

to slow evaporation.

Based on the above observations I proposed a reaction mechanism as follows:

TEOS slowly dissolved and hydrolyzed in the aqueous phase followed by

condensation to form silica nanoparticles which settled on the top of the hydrogel

phase. Simultaneously, the hydrogel at the bottom phase gradually dissolved into

the aqueous phase and disassembled into thinner fibrils. As more and more silica

nanoparticles were deposited on the top of the gel phase, a layer of silica was

formed at the interface preventing the remaining hydrogel nanofibrils from

dissolving into the aqueous phase.

2 m

(a)

237

Figure 6.5 SEM analysis on (a) the upper aqueous phase and (c) the lower hydrogel

phase. (b) EDX spectrum of (a).

Interestingly, during the sample collection at 1 month stage, the silica layer was

broken. After settling for another 1 month, the upper aqueous phase transformed to a

very weak clear gel and the lower gel phase remained clear. The unstained TEM image

of the resulting clear gel revealed a network of silicified nanofibrils (Figure 6.6a)

as evidenced by the coverage of a higher contrast layer on the hydrogel

nanofibers. The unstained hydrogel nanofibrils were hardly visible under TEM

due to very weak electron contrast (Figure 6.6b). The periodicity of the coating

was not observed here. This observation suggested that after the breakage of the

silica layer at the interface, the remaining hydrogel nanofibrils entered into the

2 m

0 2 4 6

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15000

20000

25000

Co

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ts (

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.)

Energy (keV)

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Si

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(b)

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238

upper aqueous phase and subsequently catalyzed the deposition of silica. The

observed circular features were attributed to the bubbles formed in the silica

coating.

Figure 6.6 Unstained TEM images of (a) the network of silicified hydrogel

nanofibrils that were derived from the resulting clear gel. Scale bar, 100 nm. (b)

Fmoc-Y self assembled hydrogel. Scale bar, 200 nm.

6.4 Conclusion

The silicification of peptide based self-assembled supramolecular structures was

studied employing the self-assembled Fmoc-Y and Fmoc-FY hydrogels as

templates. Both of the gels were prepared through an enzyme triggered

dephosphorylation. The presence of a high density of –OH group on the

nanofibers’ surface was found to promote silica deposition.

The

(a)

(b)

239

6.5 References

1. B. Xu et al., Enzymatic Formation of Supramolecular Hydrogels, Adv.Mater.,

2004, 16, 1440.

2. J. W. Sadownik et al., Micelle to fibre biocatalytic supramolecular

transformation of an aromatic peptide amphiphile, Chem. Comm., 2011, 47, 728.

3. A. M. Smith et al., Fmoc-Diphenylalanine Self Assembles to a Hydrogel via a

Novel Architecture Based on π-π Interlocked β-Sheets, Adv. Mater., 2008, 20, 37.

240

Chapter 7 Graphene-Inorganic hybrids

7.1 GO-Inorganic nanohybrids

7.1.1 Characterization of GO dispersion

The GO sheets were found to readily dispersible in water due to the presence of

the hydrophilic functionalities, resulting in homogenous and stable colloidal

dispersions which show no signs of agglomeration over a period of 1 month.

These dispersions possessed a brown color which is characteristic of highly

oxidized GO due to the adsorption of highly oxidative aromatic debris on their

surface 1. SEM analysis (Figure 7.1a) revealed random aggregates of crumpled

GO sheets with bright-field TEM image (Figure 7.1b) showing that these sheets

were fairly flat with folds present at both sides (indicated by arrows). The SAED

pattern (Figure 7.1c) taken from the centre of the sheet (indicated by the dashed

box) exhibited sharp diffraction spots arranged in the typical sixfold symmetry.

This leads to the conclusion that short-range order is present in the structure of

GO, contradictory to some of the previous reports claiming that GO is completely

amorphous 2. Wilson et al. 3 suggested that GO is not only comprised of fully

amorphous regions due to the presence of sp3 domains formed during oxidation

but some crystalline regions (unoxidized sp2 domains) are also present. The

diffraction peaks are labeled using Miller-Bravais (hkil) indices. It has been

previously proposed that the relative intensities of the 1-100 and 2-1-10 type

reflections can be used to identify the layer numbers of graphene 4. The intensity

profile plot along the line between the arrows shown in Figure 7.1c revealed an

intensity ratio of I0-110/I1-210 and I-1010/I-2110 ~1.3 (Figure 7.1d), indicating the

monolayer nature of the GO sheet 3,4.

Figure 7.1e showed a lower magnification TEM image of GO sheets with many

folds clearly visible (indicated by arrows). The corresponding SAED pattern

(Figure 7.1f) taken from the region marked by the dashed box revealed a

superposition of three hexagonally symmetric patterns (indicated by yellow, red

241

and blue colors) rotated by 24° and 37º respectively, which is in accordance with

the observations by Wilson3 and Wang5. This is probably due to the

misorientation between three overlapped GO sheets as a result of the fold (red

arrow) present along the left side within the dashed box.

(c)

(1-210)

(1-100)

(-1010) (-2110)

(0-110)

(-1-120)

(b)

10 μm

(a

)

0 100 200 300 400 500 600

20

40

60

80

Inte

nsity (

a.u

.)

(d)

1-210

0-110

-2110

-1010

242

Figure 7.1 (a) SEM image of aggregated GO sheets. (b) TEM image of single layer

GO sheet with folds present at both sides (indicated by arrows). Scale bar, 100 nm.

(c) Corresponding SAED pattern taken from the region marked by the dashed box

in (b). The pattern was labeled with Miller-Bravais indices. (d) Intensity profile plot

along the line between the arrows shown in (c). (e) Lower magnification TEM image

of GO sheets with the folds indicated by arrows. Scale bar, 200 nm. (f)

Corresponding SAED pattern taken from the region marked by the dashed box in

(e) showing three superimposed hexagonal patterns indicated by yellow, red and

blue colours.

7.1.2 Preparation of GO-TiO2 nanohybrids

GO-TiO2 nanohybrids were prepared from the hydrolysis of TBOT in the

presence of aqueous dispersion of GO. Two concentrations of TBOT (1.5 mM

and 3.7 mM) were studied for their effect on the morphology of the hybrids

whilst keeping the other conditions constant. Upon addition of TBOT solution

into the GO dispersion, precipitation was observed immediately. TEM analysis

showed that the reaction of GO with lower concentration of TBOT for 4 h led to

uniform dispersion of TiO2 nanoparticles with an average diameter of 16.1±2.9nm

on GO sheets (Figure 7.2a). The internal mesoporous structure of the TiO2

nanoparticles was evident by the presence of voids in the form of white dots

within individual nanoparticles (Figure 7.2b). The SAED pattern (inset) taken

from the region marked by the red dashed box revealed two sharp rings resulted

from the superposition of individual diffraction spots corresponding to different

crystallographic orientations. This is likely due to the presence of folds or

(e) (f)

Can use image J,

Double click on the

“line symbol”, change

width, 划线 ,

“Analysis-plot profile”,

copy and paste in

excel.

243

wrinkles in the GO structure. The inner ring appeared more intense than the

secondary ring, which is consistent with monolayer sheets. The absence of

diffraction patterns typical for either anatase or rutile indicates that the

as-produced TiO2 nanoparticles on GO are amorphous. Elemental analysis shown

in Figure 7.2c verified the presence of titanium in the resultant nanohybrids. After

reaction for 7 d, GO was found to be covered more densely with smaller TiO2

nanoparticles (Figure 7.2d).

Significantly larger TiO2 nanoparticles with an average diameter of 59.0 ± 6.3 nm

were observed on GO sheets at higher concentration of TBOT (Figure 7.2e). The

larger particle size is likely attributed to the increased hydrolysis rate of TBOT at

higher concentration. The SAED pattern (Figure 7.2f) taken from a fairly flat

region of the sheet (indicated by the dashed box) revealed a six-fold pattern

which is consistent with the hexagonal lattice of GO 6. Intensity profile (Figure

7.2g) plot through the (2-1-10)–(1-100)–(0-110)–(-1-120) axis showed that the

inner 1-100-type reflections are roughly 1.5 fold as intense as the outer

2-1-10-type reflections, indicating monolayer GO 3.

It is likely that the TiO2 nanoparticles are covalently bonded onto GO sheets as a

result of the condensation reaction between the hydroxyl groups of TiO2 nuclei

and the oxygen-containing functionalities (–COOH and –OH) on GO sheets 7.

However, Zhou 8 and Jiang 9 have demonstrated that TiO2 nanoparticles were

physisorbed on GO sheets in the produced nanocomposites and only after

annealing, Ti-C bond was formed 10,11.

The GO-TiO2 nanohybrids synthesized in the current study exhibited a more

uniform morphology and greater coverage than those obtained from the direct

blending of preformed TiO2 nanoparticles with GO suspension 12,13.

244

(c)

(e)

(a)

(d)

(-1-120)

(0-110)

(1-210)

(1-100) (-1010)

(2-1-10)

(f)

(b)

245

Figure 7.2 (a) TEM image of GO-TiO2 nanohybrids produced with lower TBOT

concentration for 4 h. Inset corresponds to the SAED pattern taken from the region

marked by the red dashed box. (b) A magnified image of the region shown in the

orange dashed box in (a). (c) EDX spectrum of (a). Note that Cu signal is originated

from the TEM grid. (d) TEM image of the hybrids produced with lower TBOT

concentration for 7 d. (e) TEM image of the hybrids produced with higher TBOT

concentration for 4 h. (f) Corresponding SAED pattern taken from the region

marked by the dashed box in (e) and the diffraction spots are labeled using

Miller-Bravais indices. (g) Intensity profile plot along the line between the arrows

shown in (f).

The samples produced with both TBOT concentrations for 4 h were subsequently

heated in Ar at 500 °C for 2h to induce TiO2 crystallization. It is clearly observed

in Figure 7.3(a) and (b) that for both of the samples, TiO2 nanocrystals were

uniformly distributed on the surface of GO sheets after heat treatment. For the

sample produced with lower TBOT concentration, the edges of GO sheets were

explicitly identified under the TiO2 nanocrystals (indicated by arrows in Figure

7.3a). Whilst the sample produced with higher TBOT concentration showed

crumpled GO sheets (Figure 7.3b) which may be induced by the elevated

temperature used during annealing 14.

SAED analysis of the thermally treated hybrid nanosheets revealed a pattern of

concentric rings (inset in Figure 7.3a and b), suggesting their polycrystalline

nature. Both of the diffraction patterns showed good agreement with those of GO

0 100 200 300 400 500 600 700

50

100

150

In

tensity (

a.u

.)

2-1-10

1-100

0-110

-1-120

(g)

246

(Figure 7.3c and e) and anatase TiO2 (Figure 7.3d and f), indicating the

crystallization of amorphous TiO2 into anatase and that the structure of GO is

well preserved during the heat treatment. Akhavan et al. have previously reported

that heat treatment could lead to the formation of Ti–C bonds between GO and

TiO2 11. The better linkage between the two materials will facilitate the electron

transfer and potentially enhances the photocatalytic efficiency.

(a) (b)

GO

(c) (110)

(100)

(d)

(211)

(004)

(101)

Anatase

http://pubs.acs.org/action/doSearch?action=search&author=Akhavan%2C+O&qsSearchArea=author

247

Figure 7.3 TEM images of the thermally treated nanohybrids obtained from (a) the

reaction with lower TBOT concentration for 4h and (b) the reaction with higher

TBOT concentration for 4h. The inset in (a) and (b) showed the corresponding

SAED patterns which were indexed to (c,e) GO (labeled using Miller (hkl) indices)

and (d,f) anatase TiO2 respectively. Note that the upper half in (c)-(f) showed the

experimental data while the lower half in (c) and (e) showed the diffraction pattern

of GO, and that in (d) and (f) showed the simulated diffractions for anatase

according to JCPDS 21-1272.

7.1.3 Preparation of GO-SiO2 nanohybrids

GO-SiO2 nanohybrids were produced by the in-situ growth of silica nanoparticles

on GO sheets using tetrakis (2-hydroxyethyl) orthosilicate (THEOS) as

water-soluble silica precursor. The hydroxyl groups of THEOS could condense

with the oxygenated functional groups (–COOH and –OH) on GO sheets, giving

rise to uniform seeding of silica on GO surface. An advantage of using THEOS

over the commonly used TEOS is that no organic solvents are needed owing to

the complete solubility of THEOS in water.

The morphology of the resultant product was characterized by SEM. In contrast

to pristine GO which revealed an agglomeration of highly crumpled sheets

(Figure 7.4a), the synthesized nanohybrids showed a fairly flat and layered

GO

(f)

Anatase

(211)

(004)

(101)

(200)

Anatase

(e)

(110)

(100)

GO

248

structure (indicated by arrows along the edges, Figure 7.4b) with a rather smooth

surface. This may be attributed to the silica film formed on GO sheets that makes

them more rigid. A higher magnification image (Figure 7.4c) showed the partial

separation of two hybrid sheets with the silica coating acting as spacer to prevent

the re-stacking of the sheets. The EDX spectrum for the SEM sample (Figure

7.4d) confirmed the coverage of silica on GO sheets.

Figure 7.4 SEM images of (a) highly aggregated GO sheets. (b) GO-SiO2

nanohybrids with layered structure (indicated by arrows along the edges). (c)

Higher magnification image showing the partial separation of two hybrid sheets. (d)

EDX spectrum of the sample shown in (b).

TEM analysis was further performed to investigate the attachment of SiO2 on GO

sheets. As shown in Figure 7.5a, rather flat GO sheets with no visible wrinkles

2 μm

(b)

0 1 2 3 4 5

0

5000

10000

15000

20000

25000

Co

un

ts (

a.u

.)

Energy (keV)

C

O Si

Al

(d)

10 μm

(a)

1 μm

(c)

249

were covered with a dense silica film composed of closely packed silica

nanoparticles as a result of the uniform seeding of silica on GO sheets. This

observation is consistent with the SEM image showing that the flat GO sheets

were coated with a smooth silica layer. As shown in the magnified image (Figure

7.5b), the presence of the wrinkles on the underlying GO sheets suggests that they

were extended beyond the image and hence the silica coating was rather present

on certain area of GO instead of uniformly covering the entire sheet. The SAED

pattern taken from the region marked by the dashed box confirmed the presence

of GO sheet in the synthesized hybrids.

Figure 7.5 (a) Low magnification TEM image of silica coated GO sheets. (b) A

magnified TEM image showing the ripples present on the GO sheet (indicated by

the arrow). The SAED pattern taken from the region marked by the dashed box

was labeled using Miller-Bravais indices.

(a)

0.2 μm 200 nm

(b)

(2-1-10)

(1-100)

(0-110)

(-1-120)

(-1010)

(1-210)

250

To provide further evidence for the formation of silica coating and explore the

potential of such hybrids for the preparation of ultrathin 2D nanomaterials, silica

sheets were obtained from calcining the GO-SiO2 hybrids in air at 600 °C for 2 h

to remove the supporting GO sheets. TEM image of the calcined sample showed

a highly porous intact sheet with a lateral size of ~2 μm (Figure 7.6a). In contrast

to the dense silica film formed in the hybrids, silica nanoparticles with an average

size of 14 nm were obtained after calcination. The similar dimension of the

calcined sheet to that of the silica film on GO before heat treatment confirmed its

silica nature. GO-like structures were also observed as indicated by the circles in

Figure 7.6a along with the porous silica sheet. This was further examined by

SAED (Figure 7.6b) which revealed a spot pattern (indicated by dashed circles)

suggesting the presence of remaining GO. This observation suggests the

improved thermal stability of GO as a result of the attached SiO2 nanoparticles.

Similar effect has been reported for CNT-SiO2 nanohybrids 15. Longer calcination

time may be required to fully remove GO.

Figure 7.6 (a) TEM image of porous silica sheets obtained from the calcination of

GO-SiO2 hybrids. (b) Corresponding SAED pattern taken from the sample shown

in (a).

(a) (b)

251

7.2 bwGO-Inorganic nanohybrids

Wilson et al. have recently revealed the real structure of as-produced GO which

was comprised of oxidized graphene sheets decorated by highly oxidative

aromatic debris 1. The debris acts as surfactant which stabilize the graphene

dispersion in water. It can be removed using a NaOH (aq) wash to give

base-washed GO (bwGO) which is no longer hydrophilic.

The debris present on bwGO should act as a nucleating species in a similar

manner to BA. However, the debris may not be the optimal surfactants. Therefore,

Fmoc-Trp functionalized bwGO was used as a template for the preparation of

bwGO-TiO2 nanohybrids.

7.2.1 bwGO dispersion

The dispersibility of bwGO in Fmoc-Trp aqueous solution was initially studied.

The dispersion in H2O was also prepared as a control. Both of the dispersions

were allowed to stand for 35 d to evaluate their stability. As shown in Figure 7.7a,

bwGO could be temporarily dispersed in H2O with the aid of ultrasonication and

the resulting suspension was very unstable which precipitated completely after 2d.

This is because that base wash has removed majority of highly oxidative debris

leading to strong van der Waals interaction between bwGO sheets. In contrast, a

very homogenous dispersion was obtained in the presence of Fmoc-Trp which did

not show signs of precipitation until 7 d (Figure 7.7b), suggesting the critical role

of the surfactants in maintaining the dispersion. Very slow precipitation of the

dispersion was observed over a period of 35 d.

252

(a)

0 min 1 d 2 d 7 d 35 d

(b)

0 min 1 d 2 d 7 d 35 d

Figure 7.7 Photographs of the aqueous dispersion of bwGO (a) in the absence and

(b) in the presence of Fmoc-Trp. The dispersions were allowed to stand for 35 days.

TEM analysis of bwGO deposited from the dispersion in Fmoc-Trp solution

revealed crumpled sheets measuring over 1 μm in diameter (Figure 7.8a). The

corresponding SAED pattern (Figure 7.8b) observed sharp rings composed of

resolved diffraction spots due to the presence of wrinkles or overlapping with

different sheets. The two rings were indexed to the (100) and (110) planes of

graphite. The higher diffraction intensities of the inner ring than the secondary

ring indicate the exfoliation of bwGO to single layer in Fmoc-Trp solution.

253

Figure 7.8 (a) TEM image of bwGO sheets deposited from the dispersion in

Fmoc-Trp solution. (b) The corresponding SAED pattern taken from the sample

shown in (a). The pattern was labeled using Miller (hkl) indices.

7.2.2 bwGO-TiO2 nanohybrids

7.2.2.1 Reaction in aqueous solution

The preparation of bwGO-TiO2 nanohybrids in aqueous solution starts with the

dispersion of bwGO in Fmoc-Trp solution, proceeding with mixing the dispersion

with TBOT solution. The morphology of the obtained bwGO-TiO2 nanohybrids

was characterized by TEM. As shown in Figure 7.9a, bwGO sheets were found to

be evenly coated with spherical TiO2 nanoparticles which showed a relatively

broad size distribution ranging from a few to several tens of nanometers. Fast

condensation rates in aqueous solution may cause the broad particle size

distributions 16. The corresponding SAED pattern shown in Figure 7.9b revealed

two sharp rings characteristic of graphitic structures. The diffraction intensities of

the inner ring were higher than the secondary ring, suggesting that the bwGO

sheets remained well dispersed during the coating process. It was noted that the

synthesized bwGO-TiO2 nanohybrids exhibited a more uniform and denser

coating compared with GO-TiO2 nanohybrids shown in section 7.1.2. This may

due to the higher density of adsorbed Fmoc-Trp on bwGO sheet than that of the

oxidative debris present in the as-produced GO.

(a) (b) (110)

(100)

254

Figure 7.9 TEM image of bwGO-TiO2 nanohybrids prepared in aqueous solution.

The arrows indicate the wrinkles present in bwGO sheets. (b) The corresponding

SAED pattern taken from the sample shown in (a).

The existence of TiO2 in the resultant nanohybrids was proved using Raman

spectroscopy. Since the as-produced TiO2 is amorphous which shows no Raman

peaks, the nanohybrid was annealed in argon at 600 ºC to induce crystallization of

TiO2. The Raman spectra for bwGO and anatase TiO2 were also recorded for

reference. The typical Raman spectrum for bwGO (Figure 7.10a) shows two

broad peaks, the G band at 1591 cm-1 and the defect-induced D band at around

1338 cm-1. The intense D band of bwGO indicates that the oxidation induced

disorders are still present after base wash, which is consistent with the model

proposed by Wilson et al. 1. The similar frequencies of D and G bands for the

nanohybrids (Figure 7.10c) to that for bwGO (Figure 7.10a) confirmed that its

electronic structure was largely preserved after annealing, which is likely

attributed to the stabilization by the attached anatase nanocrystals. A weak peak

observed at 144 cm-1 (Figure 7.10c) is assigned to the E1g mode of anatase

phase17, thus verifying the existence of TiO2 in the produced nanohybrids.

(b) (a)

255



prepared in aqueous solution. The spectra were taken using a 633 nm HeNe laser.

Note that the peak at around 520 cm-1 was attributed to the SiO2/Si substrate.

7.2.2.2 Reaction in EtOH

Very smooth and uniform TiO2 coating was observed on CNTs with the reaction

proceeding in EtOH (section 5.3), herein, similar procedure was employed to

prepare bwGO-TiO2 nanohybrids. It starts with the dispersion of bwGO in EtOH

followed by the addition of Fmoc-Trp solution (or H2O as a control) to stabilize

the dispersion before finally mixing with the TBOT solution.

SEM analysis of the resultant nanohybrids revealed planar sheet structures which

showed a very rough surface (Figure 7.11a), suggesting the coating of TiO2

nanoparticles. The corresponding EDX spectrum (Figure 7.11b) further verified

the coverage of TiO2 on bwGO by showing Ti, O and C element peaks.

500 1000 1500 2000 2500 3000 3500

0

100000

200000

Inte

nsity (

a.u

.)

Raman shift (cm-1)

G (a) bwGO

(b) Anatase

D

(c) Annealed bwGO-TiO2

prepared in aqueous solution

D G

E1g

E1g

256

Figure 7.11 (a) SEM image of bwGO-TiO2 nanohybrids prepared in EtOH. (b) EDX

spectrum. Pt signal is originated from Pt coating on the SEM sample to reduce

charging effect.

TEM analysis was further performed to study the effect of reaction media on the

morphology of the resultant hybrids. As shown in Figure 7.12(a) and (c),

uniformly coated bwGO sheets were observed for the samples prepared in EtOH.

The size of the TiO2 nanoparticles was too small to be clearly distinguished.

Although the precise mechanism for the formation of the hybrids is unclear, we

tentatively speculate the mechanism as follows: Upon addition into EtOH

dispersion, hydrolyzed TBOT bind onto bwGO surface through both H-bonding

interaction with the adsorbed Fmoc-Trp and condensation with the oxidized

functionalities on the sheet before heterogeneous nucleation to form TiO2 nuclei

(Figure 7.13a). While upon addition into aqueous dispersion, fast hydrolysis and

condensation rates lead to homogenous nucleation in solution, and subsequently

the nuclei bind onto bwGO surface where they grow to form TiO2 nanoparticles

(Figure 7.13b). Therefore, less uniform coating of larger particles was formed.

The role of the surfactants in the formation of the hybrids was also studied. ED

analysis of the sample prepared in the absence of Fmoc-Trp revealed ring patterns

(Figure 7.12b) with the inner ring showing a higher intensity suggesting

Ti Pt

Al

0 1 2 3 4 5

0

2500

5000

Co

un

ts (

a.u

.)

Energy (keV)

C

O

(b)

1 μm

(a)

257

randomly restacked monolayer sheets. This may due to that without the

surfactants, temporarily dispersed bwGO sheets restack prior to coating with TiO2

despite the presence of oxygenated functionalities on their surface. In contrast,

the SAED pattern for the sample prepared with the addition of Fmoc-Trp solution

(Figure 7.12d) showed diffraction spots arranged in six-fold symmetry. It’s clear

from the intensity profile (Figure 7.12e) that the inner 1-100-type reflections are

more intense than the outer 2-1-10-type reflections, with an intensity ratio of ~1.5,

indicating individually coated sheets which are stabilized by Fmoc-Trp.

(a)

(b)

(c)

(-1-120)

(2-1-10)

(0-110)

(d)

(1-100)

(1-210)

(-1010)

258

Figure 7.12 (a) TEM image of bwGO-TiO2 nanohybrids prepared in EtOH with the

addition of H2O. (b) Corresponding SAED pattern taken from the region marked

by the dashed box in (a). (c) TEM image of bwGO-TiO2 nanohybrids prepared in

EtOH with the addition of Fmoc-Trp solution. (d) Corresponding SAED pattern

taken from the region marked by the dashed box in (c). The pattern was labeled

using Miller-Bravais indices. (e) Intensity profile plot along the line between the

arrows shown in (d).

0 100 200 300 400 500 600

50

100

150

200

In

ten

sity (

a.u

.)

2-1-10

1-100 0-110

-1-120

(e)

OH

OH

In EtOH

Nucleation

bwGO-TiO2 nanohybrids

~ TiO2 nanoparticle

(a)

259

Figure 7.13 Schematic illustration of the synthesis of bwGO-TiO2 nanohybrids in (a)

EtOH and (b) aqueous solution.

Raman spectroscopy was also conducted to verify the existence of TiO2 coating

in the hybrids prepared in EtOH. The hybrids were annealed in argon at 600 ºC to

induce crystallization of TiO2 as well as to avoid the oxidation of bwGO. Raman

spectrum for the annealed sample (Figure 7.14c) showed the characteristic peak

for the E1g mode of anatase TiO2 at 149 cm-1 and the broad D and G bands which

were indexed to bwGO. This result provides further evidence for the coverage of

TiO2 on bwGO sheets.

OH

OH

OH

OH

HO

HO

OH

Growth

OH

OH

OH

HO

HO

OH

OH

OH

HO

HO

OH

OH

OH

HO

HO OH

OH

OH

HO

HO

bwGO-TiO2 nanohybrids

OH

OH

OH

HO

HO

In aqueous solution

~ TiO2 nuclei

~ TiO2 nanoparticle

(b)

260



prepared in EtOH. The spectra were taken using a 633 nm HeNe laser. Note that

the peak at around 520 cm-1 was attributed to the SiO2/Si substrate.

7.3 Exfoliated graphene-Inorganic nanohybrids

By analogy with the successful dispersion of CNTs in Fmoc-AA solution, herein,

Fmoc-Trp was used as surfactant for the direct exfoliation of graphite. The

quality of the resultant dispersion (dispersed concentration, flake size and

thickness) was assessed by Raman spectroscopy, TEM and AFM prior to the

preparation of exfoliated graphene (EG)-TiO2 nanohybrids.

7.3.1 Effect of sonication time and centrifuge speed on

the concentration of the graphene dispersion

A series of graphene dispersions were prepared via ultrasonicating graphite in

Fmoc-Trp solution with various sonication time and centrifuge speed to study

their effect on the concentration of the resultant dispersion. The photographs of

the dispersions prepared with increasing sonication time (at a constant centrifuge

500 1000 1500 2000 2500 3000

0

20000

40000

60000

80000

100000

Inte

nsity (

a.u

.)

Raman shift (cm-1)

D G (a) bwGO

(b) Anatase

(c) Annealed bwGO-TiO2

prepared in EtOH

E1g

E1g

E1g

D G

261

speed of 3000 rpm) and centrifuge speed (at a constant sonication time of 6 h)

were shown in Figure 7.15(a) and (b) respectively. All of the obtained dispersions

were homogenous and the color of the dispersion prepared with longer sonication

time and lower centrifuge speed appeared darker, indicating higher concentration

of graphene. The exact concentration of the dispersion prepared under each of the

conditions was determined using a simple filtering and weighing method and is

summarized in Table 7.1. The concentration increased from 0.016 mg/mL for 1 h

of sonication to 0.03 mg/mL for 12 h of sonication while fell from 0.070 mg/mL

after centrifugation at 500 rpm to 0.015 mg/mL after centrifugation at 6000 rpm,

which is consistent with the darkness of the dispersion in the photographs. These

values are comparable to those reported in the literature for surfactant stabilized

graphene 18,19 and are approximately 10 times lower than the reported

concentration of graphene dispersed in NMP which was prepared under the same

condition 20. Although the concentration is smaller than those achieved in

solvents, working in aqueous system brings its own advantages.

Figure 7.15 Digital images of the graphene dispersions prepared under various

conditions: (a) with increasing sonication time (at a constant centrifuge speed of

3000 rpm) (b) with increasing centrifuge speed (at a constant sonication time of 6 h).

1 h 6 h 12 h

(a)

500 rpm 3000 rpm 6000 rpm

(b)

262

Table 7.1 Measured concentrations of graphene dispersions produced with various

sonication time and centrifuge speed

In addition, the stability of the dispersion was studied by monitoring the change

in the homogeneity of the dispersion over time. Similar to the colloidal

stabilization of CNTs by Fmoc-Trp, the dispersion of graphene in Fmoc-Trp

solution was found to be fairly stable over a period of 1 week by showing only

moderate degree of sedimentation.

7.3.2 Evidence for exfoliation to graphene

7.3.2.1 Raman characterization of the exfoliated samples

Raman spectroscopy has been proved as a powerful and nondestructive tool for

capturing the electronic structures of graphene 21. Therefore, to probe the degree

of exfoliation and to evaluate the quality of the exfoliated flakes, the resultant

dispersion was deposited on SiO2/Si substrate for Raman measurement. At least

40 flakes from each of the samples were randomly chosen and measured under

laser excitation of 633 nm. The spectrum for the starting graphite powder was

also recorded for comparison. All the peaks were fitted with Lorentzians. The

Raman spectrum for the starting graphite powder (Figure 7.16a) exhibited a

single and sharp G band at ~ 1581 cm-1, a negligible D band at ~1333 cm-1, and a

2D band at ~2686 cm-1. The negligible D band suggests a nearly defect-free

structure of the starting graphite. The 2D band is composed of 2 components, 2D1

and 2D2, which are roughly 1/4 and 1/2 the height of the G peak respectively.

Raman spectra for all the exfoliated samples (Figure 7.16b-d and Figure 7.17b-d)

were in striking contrast to that for graphite by showing a considerably

pronounced D band at ~1331 cm-1, a weak shoulder peak D’ at ~1617 cm-1 and a

500 3000 6000

1 0.016

6 0.070 0.024 0.015

12 0.030

Centrifuge speed (rpm)

Sonication

time (h) Concentration (mg/mL)

263

symmetric 2D band which was force fitted to a single Lorentzian peak for

comparison purposes. The 2D band position range from 2656 to 2668 cm-1 with a

bandwidth of 60-90 cm-1 which is typical for few layer graphene (< 5 layers)

dispersed in solvents or surfactant solutions 19,20,22,23. It should be noted that due

to the aggregation of the flakes on the substrate during solvent evaporation and

the larger size of the laser spot compared with that of the flakes, the obtained

Raman spectra include contributions from all the flakes in the aggregate which

lead to the broadening and symmetry of the 2D band. The I2D/IG ratio for the

exfoliated samples is found to be ~0.5, again suggesting their few layer nature 24.

A significantly stronger D band was observed for all the exfoliated samples as

compared to the starting graphite, indicating the introduction of large amounts of

defects during the sonication process. These defects can be either basal-plane

defects or defects associated with the new edges formed as the flake size is

reduced during the extensive sonication. Coleman et al. have identified the

introduced defects being associated with the formation of new flake edges rather

than body defects by showing the time dependence of the difference between

ID/IG for exfoliated graphene versus graphite powder 19,20. This implies that the

sonication process is relatively non-destructive which yields flakes with good

quality.

A weak shoulder peak D’ is observed for all the exfoliated samples, similar to the

observation in previous study of graphene exfoliated in surfactant solution, which

is mostly attributed to the edge defects 25.

It is also noted that the G band for the exfoliated samples slightly upshifted

(1583cm-1) compared with that of graphite (1581 cm-1). Graf et al. 26 have

demonstrated that only mono- and bilayer graphene show upshifted G band

compared with that of HOPG. Ferrari et al. have also reported the upshift of the G

peak of monolayer graphene compared with that of bulk graphite which was

partially attributed to chemical doping 21. Gupta et al. 27 have shown that the G

264

band frequency exhibited an almost linear dependence on 1/n (n = number of

graphene layers) and according to their result, 3 and 4 layers are dominant in the

present study.


deposited from the dispersions prepared with (b) 1 h (c) 6 h and (d) 12 h of

sonication followed by centrifugation at 3000 rpm respectively. The spectra were

measured on SiO2/Si substrate and in all cases the excitation wavelength was 633

nm. D, G, 2D and D’ bands are indicated in the Figure. All the spectra were

normalized to have the similar G band intensity and offset for clarity.

1500 2000 2500

0

5000

10000

15000

20000

25000

30000

35000

40000

In

ten

sity (

a.u

.)

Raman shift (cm-1)

D

G 2D

D’

(a)

(b) 1 h, 3000 rpm

(c) 6 h, 3000 rpm

(d) 12 h, 3000 rpm

Graphite

265


deposited from the dispersions prepared with centrifugation at (b) 500 rpm (c) 3000

rpm and (d) 6000 rpm following 6 h of sonication respectively. The spectra were

measured on SiO2/Si substrate and in all cases the excitation wavelength was 633

nm. D, G, 2D and D’ bands are indicated in the Figure. All the spectra were

normalized to have the similar G band intensity and offset for clarity.

Approximately 40 flakes produced under each condition were analyzed to

generate the statistics that allow further investigation of the effect of sonication

time and centrifuge speed on the degree of exfoliation and flake dimensions.

Firstly, the 2D band position was considered. The statistical data for the 2D band

position under varying conditions are plotted as histograms in Figure 7.18. Figure

7.19 shows the mean 2D band position as a function of sonication time and

centrifuge speed, as calculated from the distributions in Figure 7.18. The mean

2D band position stayed roughly constant at 2662 cm-1 for all the sonication time

used (Figure 7.19a). In contrast, it shifted to lower wavenumber with increasing

centrifuge speed (Figure 7.19b) suggesting that thinner flakes remain dispersed at

higher rotation rate.

1500 2000 2500

0

5000

10000

15000

20000

25000

30000

35000

40000

Inte

nsity (

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.)

Raman shift (cm-1)

D

G 2D

(a) Graphite

(b) 6 h, 500 rpm D’

(c) 6 h, 3000 rpm

(d) 6 h, 6000 rpm

266

Figure 7.18 Histograms and normal distribution of the 2D band position for varying

sonication time and centrifuge speed.

1h, 3000 rpm

2652 2656 2660 2664 2668 2672

0

2

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6

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10

Count

Raman shift (cm-1

)

(a)

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18

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un

t

Raman shift (cm-1)

6h, 500 rpm (d)

2652 2656 2660 2664 2668 26720

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10

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un

t

Raman shift (cm-1

)

6h, 3000 rpm (b)

2652 2656 2660 2664 2668 26720

2

4

6

8

Co

un

t

Raman shift (cm-1)

6h, 6000 rpm (e)

2652 2656 2660 2664 2668 26720

2

4

6

8

10

12

14

Co

un

t

Raman shift (cm-1)

12h, 3000 rpm

(c)

267

Figure 7.19 Mean 2D band position as a function of (a) sonication time and (b)

centrifuge speed. The data for the starting graphite powder was also shown for

comparison.

The statistics of the 2D bandwidth under varying conditions was also investigated.

The histograms show that the 2D bandwidth for the exfoliated samples fall into

the range of 60-90 cm-1 (Figure 7.20). The mean 2D bandwidth is found to hardly

vary with both sonication time and centrifuge speed by displaying a relatively

constant value of 77 cm-1 (Figure 7.21).

0 5 10 15

2600

2620

2640

2660

2680

2700

2D

ba

nd

po

sitio

n

Sonication time t (h)

Graphite (a)

0 2000 4000 6000 8000

2600

2620

2640

2660

2680

2700

2D

ba

nd

po

sitio

n


Graphite (b)

40 60 80 100 1200

2

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18

20

Co

un

t

band width

1h, 3000 rpm (a)

40 60 80 100 1200

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4

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14

16

Co

un

t

band width

6h, 3000 rpm (b)

268

Figure 7.20 Histograms and normal distribution of the 2D bandwidth for varying


Figure 7.21 Mean 2D bandwidth as a function of (a) sonication time and (b)

centrifuge speed.

40 60 80 100 1200

2

4

6

8

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t

band width

(c) 12h, 3000 rpm

40 60 80 100 120 140 1600

5

10

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25

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un

t

band width

6h, 500 rpm (d)

40 60 80 100 120 140 1600

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4

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un

t

band width

6h, 6000 rpm (e)

0 5 10 15

0

40

80

120

160

2D

ba

nd

wid

th

Sonication time (h)

(a)

0 2000 4000 6000 8000

0

40

80

120

160

2D

ba

nd

wid

th


(b)

269

Ni et al. have suggested that the number of graphene layers is sensitive to the

I2D/IG ratio 28. The histograms of I2D/IG ratio for the exfoliated samples (Figure

7.22) show an average value of 0.52 apart from the sample prepared with 6 h of

sonication followed by centrifugation at 500 rpm which reveals a smaller value of

0.43. Such values are comparable to those for bi- and trilayer graphene 29, thus

verifying the few layer nature of the exfoliated flakes. As shown in Figure 7.23,

the ratio stays roughly constant with varying sonication time, while increases

steadily with centrifuge speed, suggesting that thinner flakes are dominating the

dispersion at higher centrifuge rate. This observation is in good agreement with

the previous study 26.

However, Berciaud et al. showed that the I2D/IG ratio for the suspended part of

monolayer is 4 times larger than that for the supported portion which exhibited

spatially varying doping levels arise from the interaction with the substrate 30.

Therefore, the I2D/IG ratio is not reliable in estimating the number of graphene

layers due to the non-uniform adhesion of graphene on the substrate.

0.3 0.4 0.5 0.6 0.70

1

2

3

4

5

6

7

8

Co

un

t

I2D

/IG

1h, 3000 rpm (a)

0.3 0.4 0.5 0.6 0.70

1

2

3

4

5

6

7

8

Co

un

t

I2D

/IG

6h, 500 rpm (d)

0.3 0.4 0.5 0.6 0.70

2

4

6

8

Co

un

t

I2D

/IG

6h, 3000 rpm (b)

0.3 0.4 0.5 0.6 0.70

2

4

6

8

Co

un

t

I2D

/IG

12h, 3000 rpm (c)

270

Figure 7.22 Histograms and normal distribution of I2D/IG ratio for varying


Figure 7.23 Mean I2D/IG ratio as a function of (a) sonication time and (b) centrifuge

speed.

It is also critical to study the defects formed during the sonication process. We do

this by monitoring the ID/IG ratio. It is found that the ID/IG ratio for the exfoliated

samples is larger than the typical value reported for surfactant stabilized flakes

(Figure 7.24)19.

We also compare the ID/IG ratio as a function of sonictaion time and centrifuge

speed. As shown in Figure 7.25a, the ratio is surprisingly insensitive to sonication

0.3 0.4 0.5 0.6 0.70

1

2

3

4

5

6

7

8

Co

un

t

I2D

/IG

6h, 6000 rpm (e)

0 5 10 15

0.0

0.5

1.0

I 2D/I

G

Sonication time (h)

(a)

0 2000 4000 6000 8000

0.0

0.5

1.0

I 2D/I

G


(b)

271

time but rather stabilize around 1.45. This observation contrasts with the previous

study of NMP-dispersed graphene where the ratio increase gradually with

sonication time 20. In contrast, the ratio is found to increase dramatically with

increasing centrifuge speed (Figure 7.25b), suggesting that more defects are being

introduced at higher rotation rate. Similar observation was reported for graphene

dispersed in NMP 20,23.

Since as increasing centrifuge speed, smaller flakes remain dispersed in solution

which gives more edges, we may conclude that sonication induced cutting of the

flakes is the main contribution to the increased defect levels. This observation is

consistent with the previous explanation. This is further supported by the TEM

observation (Figure 7.27) that the average size of the exfoliated flakes is smaller

than that of the laser spot, thus there will always be a large quantity of edges seen

by the beam. These act as defects to give an intense D band 31. However, we

cannot rule out a contribution from the basal plane defects which are induced

during sonication. In addition, the lack of broadening of G peak and the

considerably narrower D band compared with those observed for GO and rGO 14

confirm that the intense D band comes predominantly from the new edges formed

during sonication 32. The observation also suggests that centrifugation is a

powerful tool to differentiate graphene flakes by both size and thickness.

0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 2.20

2

4

6

8

10

Co

un

t

ID/I

G

(a) 1h, 3000 rpm

0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 2.20

2

4

6

8

10

Co

un

t

ID/I

G

6h, 3000 rpm (b)

272

Figure 7.24 Histograms and normal distribution of ID/IG ratio for varying sonication

time and centrifuge speed.

Figure 7.25 Mean ID/IG ratio as a function of (a) sonication time and (b) centrifuge

speed. The ratio for the starting graphite was also shown for comparison.

0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 2.2 2.40

1

2

3

4

5

6

7

8

Co

un

t

ID/I

G

(e) 6h, 6000 rpm

0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 2.2 2.40

2

4

6

8

10

12

14

16

18

Co

un

t

ID/I

G

6h, 500 rpm (d)

0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 2.20

2

4

6

8

10

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Co

un

t

ID/I

G

12h, 3000 rpm (c)

0 5 10 15

-1

0

1

2

I D/I

G

Sonication time (h)

(a)

Graphite

0 2000 4000 6000 8000

-1

0

1

2

I D/I

G


(b)

Graphite

273

The ID/IG ratio which reflects the dimension of the flakes was plotted against the

2D band position which reflects the thickness of the flakes for all the conditions

used. As shown in Figure 7.26, the sample prepared with 6 h of sonication

followed by centrifugation at 6000 rpm exhibits the most redshifted 2D band with

respect to that of graphite, indicating the fewest number of layers per flake while

it shows the highest ID/IG ratio suggesting the smallest size of the flakes. As the

thickness of graphene is the key factor for the property of graphene based

nanocomposites, the dispersion prepared with 6h/6000 rpm is selected for the

subsequent preparation of EG-TiO2 nanohybrids.

Figure 7.26 Plot of ID/IG ratio against 2D band position for varying sonication time

and centrifuge speed. The data for the starting graphite was also shown for

comparison. The direction of the arrow corresponds to flakes of fewer layer and

smaller size.

7.3.2.2 TEM characterization of the exfoliated samples

The lateral size of graphene flakes is very important for a number of

applications33,34. For example, large flakes with the length of a few microns or

greater are required for effective mechanical reinforcement in composites 35,36.

Therefore, TEM analysis was carried out to study the flake size distribution of the

exfoliated samples. Around 40 flakes from each of the samples were randomly

0

0.5

1

1.5

2

2.5

2655 2660 2665 2670 2675 2680 2685 2690

2D band position (cm-1)

I D/I

G

6h, 6000rpm12h, 3000rpm1h, 3000rpm6h, 3000rpm6h, 500rpmGraphite

274

chosen and measured.

Generally, well-exfoliated flakes were obtained as evidenced by the high degree

of transparency of the flakes to the electron beam (Figure 7.27). It was observed

in Figure 7.27(a) and (l) that small flakes tend to aggregate in a disordered

manner to form thicker flakes. The flakes also showed a great tendency to fold

leading to an irregular shape (Figure 7.27b-d). A few thick objects (<15%) were

occasionally observed at the lowest centrifuge speed of 500 rpm (Figure 7.27h)

which is believed to be poorly exfoliated nanographite particles. At rotation rates

of 3000 rpm and above no such very thick objects were observed in the samples.

The ED pattern of a relatively flat flake shown in Figure 7.27g revealed the

typical 6-fold symmetry (Figure 7.27i) which matches well with those observed

for NMP exfoliated graphene 32. The intensity profile plot through the

(2-1-10)–(1-100)–(0-110)–(-1-120) axis revealed an intensity ratio of I1-100/I2-1-10

~ 0.32 and I0-110/I-1-120 ~ 0.55 (Figure 7.27j), suggesting the multilayer nature of

the flakes 37 which retain the Bernal (AB) stacking of the source graphite 4.

Although this result is consistent with previously reported ratio for bilayer

graphene 4,21,32, it is hard to determine the precise number of layers for few-layer

graphene.

(a) (b)

275

(h) (g)

(c)

(f)

(i) (j)

(d)

-1010

(e)

276

Figure 7.27 Representative TEM images of graphene flakes deposited from the

dispersions prepared with (a,b) 1 h of sonication followed by centrifugation at 3000

rpm (c,d) 6 h of sonication followed by centrifugation at 3000 rpm (e,f) 12 h of

sonication followed by centrifugation at 3000 rpm (g,h) 6 h of sonication followed by

centrifugation at 500 rpm (i,j) 6 h of sonication followed by centrifugation at 6000

rpm. (k) Corresponding SAED pattern taken from the region marked by the dashed

box in (g). The pattern was labeled with Miller-Bravais indices. (l) Intensity profile

plot along the line between the arrows in (k).

The size of the flakes was determined by measuring their area using the software

AxioVision Rel. 4.8. An example of the measurement is shown in Figure 7.27d,

where a folded flake was outlined with the red line and according to the pixels of

the scale bar and that of the outlined region, the flake area is determine to be

0.1369 μm2. The statistical data for the calculated flake area under varying

conditions were summarized in Figure 7.28(a-e). The measured flake size showed

a broad distribution ranging from 0.02 to 0.89 μm2, which is lower than those

reported for solvent- 20 and surfactant-exfoliated graphene18,19. This is partly due

to the heavily folding of the flakes, leading to underestimation of the flake size.

The mean flake area appeared to be insensitive to sonication time by showing a

rather constant value around 0.13 μm2 (Figure 7.28f), which is consistent with the

Raman ID/IG ratio shown in Figure 7.25a. This observation is in contrast to that

observed for NMP-dispersed graphene where both the flake length and width

dropped off with increasing sonication time as a result of sonication-induced

(0-110) (-1010)

(1-100)

(-1-120)

(2-1-10)

(1-210)

(k)

0 100 200 300 400

50

100

150

200

250

In

ets

nity (

a.u

.)

2-1-10

1-100 0-110

-1-120

(l)

277

cutting 20. However, Lotya et al. 19 have reported a similar observation with

sodium cholate-dispersed graphene. A possible explanation is that in the presence

of the surfactants, within the first 1h of sonication, the flakes were cut to a certain

size and beyond this time sonication-induced cutting is limited possibly due to the

surfactant slippage at the graphene/surfactant/water double interface.

In contrast, the mean flake area decreases exponentially with centrifuge speed,

falling from 0.32 μm2 for sample prepared with centrifugation at 500 rpm to

0.08μm2 for that prepared with centrifugation at 6000 rpm (Figure 7.28g). This

observation agrees well with the previous work on both NMP and

surfactant-dispersed graphene 19,20. This is also consistent with the inverse ID/IG

ratio shown in Figure 7.25b, further supporting that the defects introduced during

the sonication process are predominately associated with the new edges formed

due to sonication-induced cutting, not the structural defects formed on the basal

plane.

We must also note that the statistical data derived from the TEM observation may

be biased toward higher values due to the loss of small flakes through the holes in

TEM grid.

0.0 0.1 0.2 0.3 0.4 0.50

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4

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10

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un

t

Flake area (μm2)

(a) 1h, 3000 rpm

0.0 0.1 0.2 0.3 0.4 0.50

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18

20

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Flake area (μm2)

(b) 6h, 3000 rpm

278

Figure 7.28 (a-e) Histograms and normal distribution of the flake area for varying

sonication time and centrifuge speed. (f) Mean flake area as a function of sonication

time. (g) Mean flake area as a function of centrifuge speed.

7.3.2.3 AFM characterization of the exfoliated samples

Coleman et al. have reported a method for estimating the number of graphene

0.0 0.1 0.2 0.3 0.4 0.50

1

2

3

4

5

6

7

Cou

nt

Flake area (μm2)

(e) 6h, 6000 rpm

0.0 0.1 0.2 0.3 0.4 0.50

1

2

3

4

5

6

Cou

nt

Flake area (μm2)

(c) 12h, 3000 rpm

0.0 0.2 0.4 0.6 0.8 1.00

1

2

3

4

5

6

7

Cou

nt

Flake area (μm2)

(d) 6h, 500 rpm

0 5 10 15

0.0

0.2

0.4

0.6

Fla

ke

are

a (μ

m2)

Sonication time (h)

(f)

0 2000 4000 6000 8000

0.0

0.2

0.4

0.6

0.8

Fla

ke

are

a (μ

m2)


Model

Equation

Reduced Chi-Sqr

Adj. R-Square

?$OP:F=1Model

Equation

Reduced Chi-Sqr

Adj. R-Square

?$OP:F=1

(g)

279

layers per flake by carefully counting the flake edges 18,20,32. However, it is very

difficult to identify the exact number of edges per flake based on the TEM images

obtained in the current study. It is also infeasible to identify the number of layers

based on the ED pattern, except for monolayer which shows more intense inner

diffraction spots than the outer’s. For trilayer and above, the ED pattern becomes

hardly distinguishable from that of bulk graphite.

AFM is the most commonly used tool to definitively identify the thickness of

graphene and few layer graphene flakes 18,32,38. Therefore, to provide further

evidence for the exfoliation, contact mode AFM measurement is performed. The

samples for AFM measurements were prepared by drop casting the highly diluted

dispersions onto freshly cleaved mica surface which was heated at ~110 ºC on a

hotplate to accelerate solvent evaporation.

It was very difficult to find measurable flakes for the dispersions prepared with

1h/3000 rpm, 6h/500 rpm and 6h/6000 rpm. Therefore, only those deposited from

the dispersions prepared with 12h/3000 rpm and 6h/3000 rpm were shown in

Figure 7.29. A ~1.5 μm wide flake with well defined edges was clearly visible in

Figure 7.29a. The shape of the flake appeared quite similar to those observed by

TEM. The height profiles shown in Figure 7.29b correspond to the measurements

along the white, black, blue and green lines shown in Figure 7.29a respectively. It

revealed two steps across the flake with the height of ~2.30 nm and ~1.37 nm,

suggesting partial exfoliation or folding of the flake. The height profile shown in

Figure 7.29d which corresponds to the measurement along the white line shown

in Figure 7.29c highlights a single step with the height of ~1.76 nm. Lotya et al.18

have shown the apparent height of graphene monolayers to be ~1 nm as measured

by AFM. Coleman et al. 32 have also reported the measured height for monolayer

graphene at around 1-2 nm (even up to 2.6 nm) rather than 0.34 nm which is

characteristic of individual pristine graphene sheet. This observation is probably

due to a combination of contrast issues and the presence of a residual NMP

between the monolayer and the substrate. While in the current study, both the

280

adsorbed surfactants and the attachment of smaller flakes (as evidenced by the

presence of brighter contrast region on the flake in the AFM image) may

contribute to the overestimation of the flake thickness. The measured heights for

the flakes deposited from the dispersion prepared with 12h/3000 rpm suggest the

number of layers below 3 which is in reasonable agreement with the Raman and

TEM results.

Figure 7.29(f) and (h) showed the height profiles acquired along the solid lines

shown in Figure 7.29(e) and (g) respectively. A single step from mica surface to

the flake revealed the height ranging from 3.90 to 4.63 nm, suggesting 4-8 layers

of the flakes deposited from the dispersion prepared with 6h/3000 rpm.

In addition to the exfoliated flakes, large numbers of spherical particles with

significantly brighter contrast were also observed (indicated by arrows in Figure

7.29), characterized by the height in excess of 20 nm. This may be attributed to

the aggregates of the surfactants. To test this, a control sample prepared by

depositing Fmoc-Trp solution on mica surface was imaged with AFM. As shown

in Figure 7.29i, the similar shape and thickness of the particles observed in the

control sample to those observed in the exfoliated samples confirms the presence

of surfactant aggregates.

It is necessary to remove the surfactants from the flakes in order to accurately

identify the number of layers per flake. Lotya et al.18 have reported the removal of

residual surfactants from both the substrate and the flakes by washing the

substrate with water. Therefore, in the present study, two approaches were

attempted to remove the residual Fmoc-Trp: (1) by dipping the sample in water

for 5 min followed by blowing dry with argon stream and (2) by annealing the

sample in argon at 300 ºC. However, similar AFM images were obtained after

both treatments with the aggregates of surfactants still present. As it is not

possible to fully remove the residual surfactants in the current study, AFM

measurement is very difficult to give reliable results on the accurate flake

281

thickness.

2 μm

(a)

(b) white line

0.0 0.5 1.0 1.5

16

17

18

19

20

21

Heigh

t (nm)

Distance (μm)

2.56 nm 1.29 nm

black line

0.0 0.5 1.0

16

17

18

19

20

Heigh

t (nm)

Distance (μm)

2.13 nm

blue line

0.0 0.515

16

17

18

19

20

Heigh

t (nm)

Distance (μm)

2.20 nm

green line

0.0 0.5 1.0 1.515

16

17

18

Heigh

t (nm)

Distance (μm)

1.46 nm

282

0.0 0.5 1.0

10

11

12

13

Heigh

t (nm

)

Distance (μm)

1.76 nm

(d) white line

0.0 0.5 1.0 1.5 2.06

7

8

9

10

11

12

13

14

15

16

Heigh

t (nm

)

Distance (μm)

4.01 nm

(f) white line (e)

2 μm

2 μm

(c)

283

0.0 0.534

35

36

37

38

39

Heigh

t (nm

)

Distance (μm)

3.90 nm

0.0 0.5 1.0

34

35

36

37

38

39

40

Heigh

t (nm

)

Distance (μm)

4.63 nm

(h) white line

black line

0.0 0.532

33

34

35

36

37

38

39

40

41

42

Heigh

t (nm

)

Distance (μm)

4.50 nm

blue line

0.0 0.5 1.034

35

36

37

38

39

40

41

42

43

44

Heigh

t (nm

)

Distance (μm)

3.99 nm

green line

(g)

2 μm

284

Figure 7.29 AFM characterization of the exfoliated flakes deposited from the

dispersions prepared with (a,c) 12 h of sonication followed by centrifugation at 3000

rpm and (e,g) 6 h of sonication followed by centrifugation at 3000 rpm on a 10 μm ×

10 μm mica surface. (b,d,f,h) Height profiles measured along the corresponding

lines shown in the AFM images. (i) AFM image of Fmoc-Trp solution (control)

deposited on a 20 μm × 20 μm mica surface.

7.3.3 Preparation of exfoliated graphene (EG)-TiO2

nanohybrids

7.3.3.1 Preparation of EG-TiO2 hybrids in aqueous solution

The hybrids were prepared by mixing the exfoliated graphene into a TBOT

solution. The graphene dispersions were prepared by 6 h of sonication followed

by centrifugation at either 6000 rpm or 3000 rpm. The microstructure of the

resultant products was characterized by TEM. As shown in Figure 7.30(a) and (c),

graphene flakes were uniformly embedded in the TiO2 nanoparticle aggregates

for both of the samples as evidenced by the uniform distribution of the black

stripes (marked by the red arrows) in the composites which were attributed to the

wrinkles present in the graphene. The average size of the TiO2 nanoparticles was

25.4 ± 6.5 nm. The flakes were found to protrude from the aggregates as

highlighted by the black arrows in Figure 7.30(b) and (d). The presence of

2 μm

(i)

285

graphene in the aggregates was confirmed by SAED which showed the typical

six-fold symmetry expected for graphite/graphene 39, 40 (Figure 7.30e). Due to the

strong intensity of the fuzzy ring that is originated from the aggregated

amorphous TiO2 nanoparticles, the diffraction peak corresponding to (2-1-10) was

not observed. The intensity profile plot through the (1-100)–(0-110)–(-1-120)

axis revealed an intensity ratio of I0-110/I-1-120 ~1.4 (Figure 7.30f), which is

consistent with monolayer graphene 3,4. Raman spectrum for the nanocomposites

showed three pronounced peaks at 1324, 1575 and 2646 cm-1 (Figure 7.30g)

which are characteristic for the D, G and 2D band of graphene respectively. The

2D band position suggests the presence of monolayer graphene21, which is in

good agreement with the ED analysis.

(b)

(d) (c)

(a)

50 nm

286

Figure 7.30 TEM images of the EG-TiO2 nanocomposites produced using the

dispersion prepared with (a,b) 6h/6000 rpm and (c,d) 6h/3000 rpm. (e)

Corresponding SAED pattern taken from the region marked by the dashed box in

(a). The pattern was labeled with Miller-Bravais indices. (f) Intensity profile plot

along the line between the arrows shown in (e). (g) Raman spectrum for the

nanocomposites shown in (a,b). Note that the red arrows indicate the wrinkles

present in the flakes and the black arrows indicate the edges of the flakes which

protrude from the composites. Note that the peak at around 520 cm-1 was attributed

to the SiO2/Si substrate. Scale bar, (a) 50 nm, (b) and (d) 100 nm, (c) 200 nm.

500 1000 1500 2000 2500 3000

0

20000

40000

60000

80000

100000

120000

140000

In

ten

sity (

a.u

.)

Raman shift (cm-1)

(g)

D G

2D

0 100 200 300 400

50

100

150

In

tensity (

a.u

.)

1-100 0-110

-1-120

(f)

(1-210)

(1-100)

(-1-120)

(0-110) (-1010)

(e)

287

7.3.3.2 Preparation of EG-TiO2 nanohybrids in EtOH

The reaction was performed in EtOH in order to achieve more uniform TiO2

coating on individual graphene flakes. The graphene dispersion in EtOH was

stabilized by Fmoc-Trp prior to the introduction of the TBOT solution.

As shown in Figure 7.31a, a high-aspect-ratio flake with the length of ~5 μm was

observed to be uniformly coated with considerably smaller TiO2 nanoparticles.

The magnified image (Figure 7.31b) of the protruding region shown in the dashed

box in Figure 7.31a revealed the staggered edges of five individual layers within

the flake (indicated by arrows). However, due to the dense TiO2 coverage, it is

very difficult to identify the exact number of layers. The corresponding SAED

pattern (Figure 7.31c) revealed the typical six-fold symmetry. It was noted that

the diffraction peaks within the same hexagon showed varied intensities, which

may be explained by that the surface of the hybrid was not normal to the incident

electron beam as a result of sample tilting 4. The resultant hybrid was further

characterized using Raman spectroscopy to probe the degree of exfoliation of the

flakes in the hybrids. As shown in Figure 7.31d, the Raman spectrum for the

as-produced nanohybrids exhibited three intense peaks at ~1331 (D band), 1583

(G band) and 2671 cm-1 (2D band) which are typical for multilayer graphene 21.

This result indicates the lower degree of exfoliation in EtOH as compared with

that in Fmoc-Trp solution and is consistent with the TEM observation.

(a) (b)

Staggered

edges

288

Figure 7.31 TEM images of EG-TiO2 nanohybrids prepared in EtOH. (a) Lower

magnification image of the resultant hybrid. (b) A magnified image of the region

shown in the dashed box in (a). The arrows indicate the staggered edges of

individual layers comprising the flake. (c) Corresponding SAED pattern taken from

(b). The pattern was labeled with Miller-Bravais indices. (d) Raman spectrum

acquired for the hybrids. Note that the peak at around 520 cm-1 was attributed to

the SiO2/Si substrate.

7.4 Conclusion

GO-TiO2 nanohybrids have been synthesized via the sol-gel process of TBOT in

the presence of a GO aqueous dispersion at near neutral pH and room temperature.

Relatively uniform TiO2 coverage was observed, with higher TBOT

concentration leading to significantly larger TiO2 nanoparticles. TEM and SAED

analysis confirmed the monolayer nature of the GO sheets in the hybrids.

GO-SiO2 nanohybrids was prepared by in-situ growth of silica nanoparticles on

GO using water soluble THEOS as precursors. The direct condensation between

the hydroxyl groups of THEOS and the oxygenated functional groups on GO

sheets gives rise to uniform seeding thus smooth SiO2 coating on GO. Highly

porous silica nanosheets were obtained after calcination.

The synthesis of bwGO-TiO2 nanohybrids was successfully demonstrated using

Fmoc-Trp as surfactants in both aqueous solution and EtOH. The reaction in

500 1000 1500 2000 2500 3000

0

10000

20000

30000

40000

50000

60000

Inte

nsity (

a.u

.)

Raman shift (cm-1)

G

2D

D

(d)

(2-1-10)

(1-100) (-1010)

(0-110)

(-1-120) (1-210)

(c)

289

EtOH resulted in the formation of more uniform hybrids with considerably

smaller TiO2 nanoparticles. Individually coated bwGO sheets were observed in

the presence of Fmoc-Trp, which stabilized the dispersion of bwGO during

coating process.

Finally, graphene dispersions were obtained by direct exfoliation of graphite in

Fmoc-Trp solution. The concentration of the dispersion was found to increase

with sonication time and decrease with centrifuge speed. The dispersions were

fairly stable with moderate degree of sedimentation observed over a period of 1

week. The degree of exfoliation and the quality of the exfoliated flakes was

studied by Raman spectroscopy which exhibited symmetric 2D bands. The 2D

band position, bandwidth and the I2D/IG ratio suggested that few layer graphene

(<5 layers) were dominating the dispersion. The statistical data for the 2D band

position and I2D/IG ratio suggest that thinner flakes survived the dispersion at

higher centrifuge speed. The statistical data for the ID/IG ratio and the flake size

measured based on TEM observation suggest that the defects introduced during

sonication were predominately associated with the new flake edges formed due to

sonication-induced cutting rather than the structural defects formed on the basal

plane. This implies that the sonication process is relatively non-destructive which

yields flakes with relatively good quality. The statistical data for the flake size

showed a broad distribution ranging from 0.02 to 0.89 μm2. The smaller value

compared with those reported in literatures may be attributed partly to the heavily

folding of the flakes which lead to underestimation of the flake size. AFM

analysis of the flakes revealed a typical thickness between 1.29 and 4.63 nm,

proving their multilayer nature. Both the adsorbed surfactants and the attachment

of smaller flakes may contribute to the overestimation of the flake thickness.

However, after treatment by both washing in water and annealing in argon, the

surfactants were still present which make it very difficult to accurately determine

the number of layers per flake. Together, the Raman, TEM and AFM analysis of

the exfoliated samples suggest that the dispersions predominately contain

graphene flakes with a thickness of less than 5 layers. Finally, graphene-TiO2

290

nanocomposites and hybrids were produced in aqueous solution and EtOH

respectively using the surfactant stabilized exfoliated graphene as templates. With

reaction in aqueous solution, well-exfoliated graphene flakes were uniformly

embedded in the TiO2 nanoparticle aggregates. Whilst with the reaction in EtOH,

more uniform coating of significantly smaller TiO2 nanoparticles on individual

multi-layer graphene were observed.

The noncovalent nature of the approach reported in this chapter allows the

preservation of the structures and properties of pristine graphene in the hybrids.

Furthermore, the solution-based method facilitates the deposition or film-casting

of the nanocomposites and nanohybrids on a wide range of substrates, thus may

benefit many potential applications such as in thin film technology 41.

7.5 References

1. J. P. Rourke et al., The Real Graphene Oxide Revealed: Stripping the Oxidative

Debris from the Graphene-like Sheets, Angew. Chem. Int. Ed. 2011, 50, 3173.

2. G. Wang et al., Facile Synthesis and Characterization of Graphene Nanosheets,

J. Phys. Chem. C, 2008, 112, 8192.

3. N. R. Wilson et al., Graphene Oxide: Structural Analysis and Application as a

Highly Transparent Support for Electron Microscopy, ACS Nano, 2009, 3, 2547.

4. J. C. Meyer et al., On the roughness of single- and bi-layer graphene

membranes, Solid State Commun., 2007, 143, 101.

5. Y. Wang et al., One-pot facile decoration of graphene nanosheets with Ag

nanoparticles for electrochemical oxidation of methanol in alkaline solution,

Electrochemistry Communications, 2012, 17, 63.

6. D. C. Marcano et al., Improved Synthesis of Graphene Oxide, ACS Nano, 2010,

4, 4806.

7. Y. Tang et al., Incorporation of Graphenes in Nanostructured

TiO2 Films via Molecular Grafting for Dye-Sensitized Solar Cell Application,

ACS Nano, 2010, 4, 3482.

8. K. Zhou et al., Preparation of graphene–TiO2 composites with enhanced

photocatalytic activity, New J. Chem., 2011, 35, 353.

9. G. Jiang et al., TiO2 nanoparticles assembled on graphene oxide nanosheets

with high photocatalytic activity for removal of pollutants, Carbon, 2011,

492693.

10. Z. Hu et al., Visible light driven photodynamic anticancer activity of graphene

oxide/TiO2 hybrid, Carbon, 2012, 50, 994.

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11. O. Akhavan et al., Photocatalytic Reduction of Graphene Oxide Nanosheets

on TiO2 Thin Film for Photoinactivation of Bacteria in Solar Light Irradiation, J.

Phys. Chem. C, 2009, 113, 20214.

12. W. L. Zhang et al., Fast and facile fabrication of

a graphene oxide/titania nanocomposite and its electro-responsive characteristics,

Chem.Commun., 2011, 47, 12286.






Chem.Commun., 2002, 34.

16. K. Woan et al., Photocatalytic Carbon-Nanotube–TiO2 Composites, Adv.

Mater., 2009, 21, 2233.

17. D. Eder et al., Carbon–Inorganic Hybrid Materials: The

Carbon-Nanotube/TiO2 Interface, Adv. Mater., 2008, 20, 1787.




dispersions, ACS Nano, 2010, 4, 3155.


2010, 6, 864.

21. A. C. Ferrari et al., Raman Spectrum of Graphene and Graphene Layers, Phys.

Rev. Lett., 2006, 97, 187401.

22. R. J Smith et al., The importance of repulsive potential barriers for the

dispersion of graphene using surfactants, New Journal of Physics, 2010, 12,

125008.

23. U. Khan et al., Size selection of dispersed, exfoliated graphene flakes by

controlled centrifugation, Carbon, 2012, 50,470.

24. Y. Y. Wang et al., Raman Studies of Monolayer Graphene: The Substrate

Effect, J. Phys. Chem. C, 2008, 112, 10637.

25. L.M. Malard et al., Raman spectroscopy in graphene, Phys. Rev., 2009, 473,

51.

26. D. Graf et al., Spatially Resolved Raman Spectroscopy of Single- and

Few-Layer Graphene, Nano Lett., 2007, 7, 238.

27. A. Gupta et al., Raman Scattering from High-Frequency Phonons in

Supported n-Graphene Layer Films, Nano Lett., 2006, 6, 2667.

28. Z. H. Ni et al., Graphene Thickness Determination Using Reflection and

Contrast Spectroscopy, Nano Lett., 2007, 7, 2758.

29. K. Wang et al., Ni induced few-layer graphene growth at low temperature by

pulsed laser deposition, AIP Advances, 2011, 022141.

30. S. Berciaud et al., Probing the Intrinsic Properties of Exfoliated Graphene:

Raman Spectroscopy of Free-Standing Monolayers, Nano Lett., 2009, 9, 346.

31. A. C. Ferrari et al., Raman spectroscopy of graphene and graphite: Disorder,

electron–phonon coupling, doping and nonadiabatic effects, Solid State

Communications, 2007, 14347.

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32. Y. Hernandez et al., High-yield production of graphene by liquid-phase

exfoliation of graphite, Nanotechnology, 2008, 3, 563.

33. P. N. Nirmalraj et al., Nanoscale Mapping of Electrical Resistivity and

Connectivity in Graphene Strips and Networks, Nano Lett., 2011, 11, 16.

34. P. E. Lyons et al., The relationship between network morphology and

conductivity in nanotube films, J. Appl. Phys., 2008, 104, 044302.

35. L. Gong et al., Interfacial Stress Transfer in a Graphene Monolayer

Nanocomposite, Adv. Mater., 2010, 22, 2694.

36. P. May et al., Approaching the theoretical limit for

reinforcing polymers with graphene, J. Mater. Chem., 2012, 22, 1278.

37. S. Horiuchi et al., Carbon Nanofilm with a New Structure and Property, Jpn J.

Appl. Phys. Lett., 2003, 42, L1073.

38. S. Lee et al., Wafer Scale Homogeneous Bilayer Graphene Films by Chemical

Vapor Deposition, Nano Lett., 2010, 10, 4702.

39. J. C. Meyer et al., The structure of suspended graphene sheets, Nature, 2007,

446, 60.

40. J. C. Meyer et al., On the roughness of single- and bi-layer graphene

membranes, Solid State Commun., 2007, 143, 101.



6, 458.

293

CHAPTER 8 General conclusions and future work

8.1 General conclusions

Aligned CNT arrays were successfully grown on oxidized silicon substrates

through an injection CVD method using ferrocene as catalyst and toluene as

carbon source. The non-covalent functionalization of CNTs was studied through

the adsorption of a library of aromatic Fmoc-AAs on both aligned CNT arrays

and randomly aligned CNT networks. The adsorption kinetics and equilibrium of

single Fmoc-AA species on the surface were measured using UV-Vis

spectroscopy. Among the Fmoc-AAs studied, Fmoc-Trp was found to have the

best affinity for CNTs by showing the highest equilibrium loading and initial

adsorption rate, whilst the non-aromatic control Fmoc-Gly showed the lowest

equilibrium loading. These results confirmed that the Fmoc group was an

efficient anchor, which could be improved by the addition of an aromatic amino

acid. The effect of the aromatic ligands on the affinity for CNTs was also

investigated employing BA as surfactant. The poorer affinity of BA for CNTs

compared with the Fmoc-AAs was possibly attributed to the lower degree of

aromaticity of the benzyl ring compared with that of the fluorenyl ring of the

Fmoc group. The fully reversible nature of the binding process was demonstrated

via the desorption of the Fmoc-AAs from CNTs’ surface in excess of water. The

equilibrium data were found to well follow the Freundlich isotherm model. Both

the higher binding energy and flat conformations of Fmoc-Trp on CNTs’ surface

contributed to the higher adsorption capacity compared with Fmoc-Gly. The

competitive binding from the library of Fmoc-AAs on graphite was developed

which efficiently identified Fmoc-Trp as the strongest binding candidate, leading

to a significantly different binding behavior compared with the individual

adsorption experiments. This approach provides an efficient way to screen a wide

range of binding candidates with similar binding energies simultaneously for the

strongest binder. A switchable surface chemistry was demonstrated to verify the

hypothesis that the Fmoc-AAs with a higher binding energy could displace those

294

with a lower binding energy from the surface. This study on the dynamic

interaction between these aromatic amino acid derivatives and CNTs not only

provides a step forward for their bioapplication where an effective and

well-studied interface is required but also pave the way towards the subsequent

utilization of these functionalized CNTs as templates for the production of hybrid

materials.

An in-situ sol-gel method was employed for the synthesis of CNT-SiO2 and

CNT-TiO2 nanohybrids using the Fmoc-AA functionalized CNTs as templates at

low temperature and neutral pH. Both SEM and TEM observation confirm the

formation of uniform SiO2 coating from TEOS on individual CNTs in the

presence of Fmoc-Trp and Fmoc-His. Fmoc-Tyr functionalized CNTs resulted in

a mixture of partially coated and uncoated CNTs due to the poorer affinity of

Fmoc-Tyr for CNTs as well as its weaker H-bonding interaction with the

hydrolyzed TEOS. A similar dependence of the degree of coating on the

Fmoc-AA structure was observed for coating CNTs with TiO2 from TBOT,

except for the case of Fmoc-Tyr which led to uniform coating of TiO2. The

morphology of the hybrids was found to be highly dependent on the CNT to

TBOT ratio and the modifier to CNT ratio. The coating of CNTs functionalized

with BA, as reported by Eder et al., was also studied. It was found that the TiO2

coating obtained on these BA functionalized CNTs was not as uniform as on the

Fmoc-AA functionalized nanotubes. This difference in coating was attributed to

the weaker π–π stacking interactions between BA and CNTs as a result of the

lower degree of aromaticity of the benzyl ring. The formation mechanism of the

SiO2 and TiO2 coating on the Fmoc-AAs functionalized CNTs was proposed. The

surface modifiers were believed to play a dual role: (1) To stabilize CNT

dispersion by acting as electrostatic surfactants. (2) To render the templates’

surface with the functionalities that catalyze silica and titania deposition.

Anatase TiO2 NTs were obtained after calcination of the CNT-TiO2 nanohybrids

as confirmed by XRD and ED analysis. Both the inner diameter and wall

295

thickness of the synthesized TiO2 NTs were found to be controlled by varying the

dimension of CNT templates and the ratio of CNT to TBOT. Factors including

heating temperature, pre-treatment and ramp rate were found to affect the phase

transformation from anatase to rutile. A simple route toward the production of

TiO2 NT arrays was also demonstrated using the CVD grown vertically aligned

CNT arrays as templates in the presence of the Fmoc-AAs. This method avoids

extended sonication which causes shortening of the tubes, therefore TiO2 NTs of

up to ~300 μm long were obtained which may benefit their application in

dye-sensitized solar cells. SEM analysis confirmed the good alignment of the

NTs in the arrays although the vertical alignment with respect to the substrate was

lost due to the wet-chemical techinique used. Such structures may be useful for

the development of electronic units.

Two populations of CNTs, one functionalized with Fmoc-His and the other with

Fmoc-Tyr which expose imidazole (hydrogen bonding) and hydroxyl

(nucleophilic) functionalities respectively were combined for the biomimetic

catalytic synthesis of silica and titania. SEM characterization revealed two

distinct assembled morphologies; bundled fibers and spherical aggregates which

were capable of catalyzing SiO2 and TiO2 deposition.

GO-TiO2 and GO-SiO2 nanohybrids have been synthesized via the sol-gel

process of TBOT and THEOS in the presence of GO aqueous dispersion at near

neutral pH and room temperature. Relatively uniform TiO2 coverage on

monolayer GO sheets was observed, with higher TBOT concentration leading to

significantly larger TiO2 nanoparticles. The direct condensation between the

hydroxyl groups of THEOS and the oxygenated functional groups on GO sheets

gives rise to uniform seeding thus smooth SiO2 coating on GO.

The synthesis of bwGO-TiO2 nanohybrids was successfully demonstrated using

Fmoc-Trp as surfactants in both aqueous solution and EtOH. The morphology of

the hybrids was found to be dependant on the reaction medium. The reaction in

296

EtOH resulted in the formation of more uniform hybrids with considerably

smaller TiO2 nanoparticles. Individually coated bwGO sheets were observed in

the presence of Fmoc-Trp, which stabilized the dispersion of bwGO during

coating process.

Graphene dispersions were obtained by direct exfoliation of graphite in Fmoc-Trp

solution. The concentration of the dispersion was found to increase with

sonication time and decrease with centrifuge speed. The dispersions were fairly

stable with moderate degree of sedimentation observed over a period of 1 week.

The degree of exfoliation and the quality of the exfoliated flakes was studied by

Raman spectroscopy which exhibited symmetric 2D bands. The 2D band position,

bandwidth and the I2D/IG ratio suggested that few layer graphene (<5 layers) were

dominating the dispersion. The statistical data for the ID/IG ratio and the flake size

measured based on TEM observation suggest that the defects introduced during

sonication were predominately associated with the new flake edges formed due to

sonication-induced cutting rather than the structural defects formed on the basal

plane. This implies that the sonication process is relatively non-destructive which

yields flakes with relatively good quality. The statistical data for the flake size

showed a broad distribution ranging from 0.02 to 0.89 μm2. AFM analysis of the

flakes revealed a typical thickness between 1.29 and 4.63 nm, proving their

multilayer nature. Together, the Raman, TEM and AFM analysis of the exfoliated

samples suggest that the dispersions predominately contain graphene flakes with

a thickness of less than 5 layers. Subsequently, graphene-TiO2 nanocomposites

and hybrids were produced in aqueous solution and EtOH respectively employing

the surfactant stabilized exfoliated graphene as templates. With reaction in

aqueous solution, well-exfoliated graphene flakes were uniformly embedded in

the TiO2 nanoparticle aggregates. Whilst with the reaction in EtOH, more

uniform coating of significantly smaller TiO2 nanoparticles on individual

multi-layer graphene were observed.

The noncovalent approach reported in this study allows the preservation of the

297

structures and properties of pristine CNTs and graphene in the hybrids.

Furthermore, the solution-based method facilitates the deposition or film-casting

of the nanocomposites and nanohybrids on a wide range of substrates, thus may

benefit potential applications such as in thin film technology 1.

Finally, a preliminary study on the silicification of Fmoc-Y and Fmoc-FY

self-assembled hydrogels was conducted. Both of the gels were successfully

prepared under physiological conditions through an enzyme triggered

dephosphorylation and found to template silica deposition. The presence of a high

density of -OH groups on the nanofibers’ surface was believed to promote the

deposition of silica.

8.2 Recommendation for future work

It would be interesting to investigate the effect of the number of binding sites on

the interaction with CNTs using a dynamic combinational library of AAs. For

instance, to compare the adsorption behavior of Fmoc-Trp with that of

Fmoc-Trp-Trp-Gly and Fmoc-Trp-Trp-Trp-Gly on CNTs.

The concentration of the graphene dispersion achieved in the present study is still

too low for many of the applications. It would be interesting to employ a 2 cycle

dispersion strategy which has been reported previously to give extremely high

concentration of (up to 28 mg/mL) dispersed graphene in NMP 2.

The electrical property of the produced aligned TiO2 NT arrays could be

evaluated for their potential as electrode materials in dye-sensitized solar cells

and for water splitting.

8.3 References 1. S. De et al., Flexible, transparent, conducting films of randomly stacked


298

6, 458.

2. U. Khan et al., Solvent-exfoliated graphene at extremely high concentration,

Langmuir, 2011, 27, 9077.

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